Adhesion - Current Research and Application

Adhesion Edited by Wulff Possart Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH ...

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Adhesion Edited by Wulff Possart

Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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Adhesion Current Research and Applications

Edited by Wulff Possart

The Editor Prof. Dr. Wulff Possart Univ. des Saarlandes, LS Polymere u. Thermodynamik Postfach 151150 66041 Saarbrücken Germany

n All books published by Wiley-VCH are carefully produced. Nevertheless, authors, editors, and publisher do not warrant the information contained in these books, including this book, to be free of errors. Readers are advised to keep in mind that statements, data, illustrations, procedural details or other items may inadvertently be inaccurate.

Library of Congress Card No.: applied for British Library Cataloguing-in-Publication Data: A catalogue record for this book is available from the British Library. Bibliographic information published by Die Deutsche Bibliothek Die Deutsche Bibliothek lists this publication in the Deutsche Nationalbibliografie; detailed bibliographic data is available in the Internet at © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim, Germany All rights reserved (including those of translation in other languages). No part of this book may be reproduced in any form – by photoprinting, microfilm, or any other means – nor transmitted or translated into a machine language without written permission from the publishers. Registered names, trademarks, etc. used in this book, even when not specifically marked as such, are not to be considered unprotected by law. Typsetting K+V Fotosatz GmbH, Beerfelden Printing betz-druck GmbH, Darmstadt Binding J. Schäffer GmbH, Grünstadt Printed in the Federal Republic of Germany Printed on acid-free paper ISBN-13: 978-3-527-31263-4 ISBN-10: 3-527-31263-3

V

Preface Adhesion is a term with many facets and different meanings to people but we always imply that there is a joint of different materials that can resist mechanical loading. In fundamental research, fundamental adhesion first of all summarizes all elementary processes at the interface that make two solid materials stick together. That interface is formed when one of the materials in a liquid state (the adhesive) is brought into contact with the other solid one (the adherend) and then solidifies. Close contact between the wetting adhesive molecules and the solid surface give rise to attractive forces known as physical intermolecular interactions (Van der Waals forces), as chemical bonds of any form or as the electric double layer which is created by mobile charges that interdiffuse through the contact due to the initial difference of electrochemical potential between adhesive and adherend. These three types of interactions are in mind when we speak about fundamental adhesion processes. In most cases, adhesives are based on polymers, either as part of the adhesive composition or polymerizing inside the adhesive after it has been brought in contact with the adherend. In compatible polymer-polymer composites, the macromolecules interdiffuse through the contact and create an interpenetration layer capable to bear mechanical load. Therefore, that interdiffusion process is considered as a fourth fundamental adhesion process. All these adhesion processes have been subject to intense research for many decades now. Nevertheless, we are still far from understanding them in such detail that we could transfer our knowledge straight to the sector of engineering and application. This current state has to do with the discovery that the action of fundamental adhesion forces is not restricted to the interface. They not only fix some layer of adhesive molecules on the surface of the adherend but they can exert strong influence on the formation of chemical and morphological structure as well as on molecular mobility in the adjacent region of the adhesive during solidification. Hence an interphase is formed. These interphases depend on the combination of adhesive and adherend surface and on the process of contact formation as well. Due to its distinct structure, the interphase possesses properties that can be much different from the behavior of the bulk adhesive. Compared to the complexity of the problem, we just start understanding what is behind interphases. Looking at the field of applied research, engineering and technology, adhesion might be related first of all with adhesive joints for most different materials and Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

VI

Preface

most different purposes everywhere in industry today. Moreover, adhesion plays a crucial role in any kind of material composite, in lacquers, paints, coatings as well as on membranes and thin films. Self-cleaning surfaces, friction and wear are other issues where adhesion is involved. Entering biology and medicine, we find most complex processes of cell adhesion and for protein adhesion on surfaces which are key problems in biocompatibility of materials for prosthetics, stents, artificial organs, suture material and surgical glues, just to mention some examples. That almost endless list of applications with the vast variability of demands is creating new questions to researchers and engineers every day. At the first glance, aspects of practical adhesion like mechanical deformation behavior, strength and durability under duty conditions are in the focus of interest. However, we have been learning from the materials testing conducted for many decades that understanding of the technical performance of adhesive joints and material compounds needs the knowledge about the fundamental processes that influence and change structure and chemistry in the adhesive. This technical challenge creates broad and continuing research efforts. It is an important task to make the results accessible to everyone interested in adhesion both from the scientific and the practical point of view. This book provides a collection of 34 contributions on many aspects of adhesion research and application written by scientists from around the globe. Their texts are based on lectures held at the 7th European Conference on Adhesion (EURADH) in Freiburg (Germany) in September 2004 which was organized by the German DECHEMA (Society for Chemical Engineering and Biotechnology) in cooperation with the French Section Française de l’Adhésion of the Société Francaise du Vide and the British Society for Adhesion and Adhesives. On that basis, the book intends to bridge current issues, aspects and interests from fundamental research to technical applications. In seven chapters, the reader will find an arrangement of latest results on fundamental aspects of adhesion, on adhesion in biology, on chemistry for adhesive formulation, on surface chemistry and pretreatment of adherends, on mechanical issues, non-destructive testing and durability of adhesive joints, and on advanced technical applications of adhesive joints. Prominent scientists review the current state of knowledge about the role of chemical bonds in adhesion, about new resins and nanocomposites for adhesives, and about the role of macromolecular architecture for the properties of hot melt and pressure sensitive adhesives. Thus, insight into detailed results and broader overviews as well can be gained from the book. Finally, the editor would like to express his sincere thanks to all authors who have contributed to that book. We all thank Dr. Hubert Pelc of Wiley-VCH and all other staff involved for their helpful assistance in preparing this book, and we hope that our readers from the scientific community and from application and engineering will benefit from the result of our work. Saarbrücken, Summer 2005

Wulff Possart

VII

Contents Preface

V

LIst of Contributors 1

1.1 1.2 1.3 1.4

2

2.1 2.2 2.3 2.4

3

3.1 3.2

XXIII

The Interfacial Chemistry of Adhesion: Novel Routes to the Holy Grail? 1 J. F. Watts Abstract 1 Introduction 1 Development of a Model Interphase 3 The Buried Interface 8 Conclusion 15 Acknowledgments 15 References 15 Modeling Fundamental Aspects of the Surface Chemistry of Oxides and their Interactions with Coupling Agents 17 P. Schiffels, M. Amkreutz, A. T. Blumenau, T. Krüger, B. Schneider, T. Frauenheim, and O.-D. Hennemann Abstract 17 Introduction: Atomistic Simulations in Adhesion 17 Prediction of Surface Properties: Ideal Reconstructions on a-SiO2 (0001) 19 Organic Components of the Adhesive and Substrate-Adhesive Interaction 23 Conclusion and Outlook 29 References 30 Adhesion at the Nanoscale: an Approach by AFM 33 M. Brogly, O. Noel, G. Castelein, and J. Schultz Abstract 33 Introduction 34 Materials and Methods 34


VIII

Contents

3.2.1 3.2.2 3.2.3 3.2.4 3.3 3.3.1 3.3.1.1 3.3.1.2 3.3.1.3 3.3.1.3.1 3.3.1.3.2 3.3.1.3.3 3.3.1.3.4 3.3.2 3.3.3 3.3.3.1 3.3.3.2 3.4

4

4.1 4.2 4.2.1 4.2.2 4.2.2.1 4.2.2.2 4.2.3 4.3 4.3.1 4.3.2 4.3.3 4.4

Preparation of Oxidized Silica Surface 35 Grafting of Functionalized SAMs onto Silicon Wafer 35 Crosslinking and Functionalization of PDMS Networks 35 Characterization of the SAMs 36 Results and Discussion 37 Force–Distance Curve Measurements and AFM Calibration 37 Force–Distance Curve Features 37 The DD Curve (Contact Mode) 37 AFM Calibration 38 Determination of the Spring Constant of the Cantilever 38 Nonlinearity of the Quadrant of Photodiodes 38 Scan Rate of the Cantilever 38 Systematic Check 39 Force–Distance Curves on Rigid Systems of Controlled Surface Chemistry 39 Force–Distance Measurements on Polymers 40 Force–Indentation Measurements on Polymers 40 Force–Indentation Curves on Systems of Controlled Surface Chemistry and Controlled Mechanical Properties 42 Conclusion 45 References 45 Organization of PCL-b-PMMA Diblock Thin Films: Relationship to the Adsorption Substrate Chemistry 47 T. Elzein, M. Brogly, and J. Schultz Abstract 47 Introduction 47 Materials and Methods 48 PCL-b-PMMA Diblocks 48 Infrared Spectroscopy 49 Transmission 49 Polarization-Modulation Infrared Reflection–Absorption Spectroscopy (PM-IRRAS) 49 Atomic Force Microscopy (AFM) 50 Results and Discussion 50 PCL-b-PMMA Bulk Characterization 50 PCL-b-PMMA Thin Films on OH-Functionalized Gold Substrates 51 PCL-b-PMMA Thin Films on Gold Substrates 55 Conclusion 56 References 57

Contents

5

5.1 5.2 5.3 5.3.1 5.3.2 5.3.3 5.4

6

6.1 6.2 6.2.1 6.2.2 6.2.3 6.2.4 6.3 6.3.1 6.3.2 6.3.3 6.4

7

7.1 7.2 7.2.1 7.2.2 7.2.3 7.2.4 7.2.5 7.2.6 7.2.7

Adhesion and Friction Properties of Elastomers at Macroscopic and Nanoscopic Scales 59 S. Bistac and A. Galliano Abstract 59 Introduction 59 Materials and Methods 60 Results and Discussion 62 Adherence Energy 62 Macroscale Friction 64 Nanoscale Friction and Adhesion 65 Conclusion 68 References 69 Chemical Structure Formation and Morphology in Ultrathin Polyurethane Films on Metals 71 C. Wehlack and W. Possart Abstract 71 Introduction 71 Materials and Methods 72 Sample Preparation 72 Experimental Characterization 74 IR Spectra Calculation 74 IR Band Assignment 75 Results and Discussion 76 Curing at Room Temperature 76 Morphology of Thin Films 79 Chemical Structure of Cured Films 80 Conclusion 85 Acknowledgments 86 References 87 Properties of the Interphase Epoxy–Amine/Metal: Influences from the Nature of the Amine and the Metal 89 M. Aufray and A. A. Roche Abstract 89 Introduction 89 Materials and Methods 90 Materials 90 Thermal Analysis (DSC) 91 Micro-Infrared Spectroscopy (l-FTIR) 91 Fourier Transform Near-Infrared Spectroscopy (FT-NIR) 92 Inductively Coupled Plasma Spectroscopy (ICP) 92 X-Ray Diffraction (XRD) 92 Polarized Optical Microscopy (POM) Coupled with a Hot Stage Apparatus 92

IX

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Contents

7.3 7.3.1 7.3.2 7.3.3 7.3.4 7.3.5 7.3.6 7.4

Results and Discussion 93 Interphase Formation Mechanisms 93 Formation of New Networks 94 Crystallization of “Modified” IPDA 94 Modification of Mechanical Properties 95 Comparison of Coatings and Metal–Bulk Interphases Influence of the Stoichiometric Ratio 100 Conclusion 101 Acknowledgments 102 References 102

8

Mapping Epoxy Interphases 103 M. Munz, J. Chung, and G. Kalinka Abstract 103 Introduction 104 Stiffness Mapping by Indentation Techniques 106 SFM-Based Stiffness Mapping in Force Modulation Microscopy (FMM) Mode 106 Depth-Sensing Micro-indentation (DSI) 108 Some Fundamental Aspects of Interphase Mapping by Indentation Techniques 110 Artifacts Induced by Topography 110 Artifacts Induced by the Extent of the Stress Field Beneath the Indenter 114 Two Cases of Mapped Epoxy Interphases 116 The Cu/Epoxy Interphase 116 The PVP/Epoxy Interphase 118 Conclusion 121 Acknowledgments 122 References 122

8.1 8.2 8.2.1 8.2.2 8.3 8.3.1 8.3.2 8.4 8.4.1 8.4.2 8.5

9

9.1 9.2 9.2.1 9.2.2 9.2.3 9.2.4 9.3

97

Mechanical Interphases in Epoxies as seen by Nondestructive High-Performance Brillouin Microscopy 125 J. K. Krüger, U. Müller, R. Bactavatchalou, D. Liebschner, M. Sander, W. Possart, C. Wehlack, J. Baller, and D. Rouxel Abstract 125 Introduction 125 Brillouin Spectroscopy on Thermal Phonons and Other Elementary Excitations 126 An Introduction to the Physics of Classical Brillouin Spectroscopy 126 The Kinematic View of Brillouin Spectroscopy 129 Scattering Geometries and Other Pitfalls 129 Brillouin Microscopy 132 Mechanical Interphases at Polymer–Substrate Interfaces 134

Contents

9.3.1 9.3.2 9.3.3 9.3.3.1 9.3.3.2 9.3.3.3 9.4

The Polymer Model System 134 Epoxy/Silicone Rubber Interphase 134 Epoxy/Metal Interphases 136 Technical Bulk Metals: Cu, Al 137 Thin Evaporated Metal Substrates: Al, Cu, Au, Mg 138 Discussion 141 Conclusion 142 Acknowledgments 142 References 142

10

Structure Formation in Barnacle Adhesive 143 M. Wiegemann Abstract 143 Introduction 143 Barnacles 144 General Aspects of Barnacle Settlement 144 Biochemical Characterization of Barnacle Cement 145 Substrate-Specific Formation of Barnacle Adhesive 146 Substrate-Specific Morphology of Barnacle Base 147 Phenomenological Approach to Adhesive Structure Formation and Morphology Changes 148 Homologous (?) Structure Formation of Biological Adherates on Hydrophobic Surfaces 150 Theoretical Colloid Approach to Structure Formation in Barnacle Adhesive 152 Conclusions 154 Acknowledgments 154 References 154

10.1 10.2 10.2.1 10.2.2 10.2.3 10.2.4 10.2.5 10.3 10.4 10.5

11

11.1 11.2 11.2.1 11.2.2 11.2.3 11.2.4 11.2.5 11.2.6 11.2.7 11.2.8

Adhesion Molecule-Modified Cardiovascular Prostheses: Characterization of Cellular Adhesion in a Cell Culture Model and by Cellular Force Spectroscopy 157 U. Bakowsky, C. Ehrhardt, C. Loehbach, P. Li, C. Kneuer, D. Jahn, D. Hoekstra, and C.-M. Lehr Abstract 157 Introduction 158 Materials and Methods 160 Chemicals for the Modification 160 Implant Materials 160 Modification of the PTFE Surface 160 Scanning Force Microscopy 162 Fourier Transform Infrared Spectroscopy 163 Environmental Scanning Electron Microscopy 163 Confocal Laser Scanning Microscopy (CLSM) 163 Isolation and Culture of HUVECs 164

XI

XII

Contents

11.2.9 11.3 11.3.1 11.3.2 11.3.2.1 11.3.2.2 11.3.3 11.4

Endothelialization of PTFE Films 164 Results and Discussion 165 Wet-Chemical Modification of PTFE Polymer Film Cell Adhesion Experiments 166 Adhesion and Cultivation in Static Culture 166 Perfusion Experiments 166 Cell Adhesion Force Measurements 167 Conclusion 169 Acknowledgments 170 References 171

12

Surface Engineering by Coating of Hydrophilic Layers: Bioadhesion and Biocontamination 175 G. Legeay and F. Poncin-Epaillard Abstract 175 Introduction 175 The Need for Bioadhesion of Biomaterials 175 Mechanism of Bioadhesion 176 Surface Engineering 177 Surface Preparation 177 Surface Sterilization 178 Results and Discussion 178 Hydrophobic Cold Plasma Treated Surfaces in Ophthalmology 178 Hydrophilic Cold Plasma Treated Surfaces Based on Polyvinylpyrrolidone (PVP) or Natural Derivative Coatings 179 Grafting of Monomer onto Plasma-Pretreated Surfaces 180 Coating with Commercial Native or Synthetic Polymers 181 Examples 183 With Different Biomolecules, i.e., Proteins 184 Implantation (ex in vivo) 184 In vivo Implantation 185 Conclusion 186 References 187

12.1 12.1.1 12.1.2 12.2 12.2.1 12.2.2 12.3 12.3.1 12.3.2 12.3.2.1 12.3.2.2 12.3.3 12.3.3.1 12.3.3.2 12.3.3.3 12.4

13

13.1 13.2 13.2.1 13.2.2 13.2.3 13.3 13.3.1 13.3.2

165

New Resins and Nanosystems for High-Performance Adhesives 189 R. Mülhaupt Abstract 189 Introduction 190 Tailor-Made Polymers and Properties on Demand 190 Controlled Polymerization and Catalysis 191 Functional Polymers from the Life Sciences 192 Reactive Extrusion and Isocyanate-Free Polyurethane Chemistry 193 Nanosystems 194 The Nano Challenge 194 Nanophase Separation 196

Contents

13.3.3 13.3.4 13.4

Nanomolecules as Molecular Nanoparticles 198 POSS and Nanocomposites 200 Conclusion 201 Acknowledgments 202 References 202

14

Influence of Proton Donors on the Cationic Polymerization of Epoxides 205 A. Hartwig, K. Koschek, and A. Lühring Abstract 205 Introduction 206 Initiators for the Cationic Polymerization of Epoxides 207 Influence of Moisture on the Polymerization Kinetics 209 Modification of the Polymerization Behavior by the Addition of Alcohols 212 Conclusion 215 Acknowledgments 215 References 215

14.1 14.2 14.3 14.4 14.5

15

15.1 15.2 15.2.1 15.2.2 15.3 15.4

16

16.1 16.2 16.2.1 16.2.2 16.2.3 16.3 16.3.1 16.3.2 16.3.3 16.4

Novel Adhesion Promoters Based on Hyperbranched Polymers A. Buchman, H. Dodiuk-Kenig, T. Brand, Z. Gold, and S. Kenig Abstract 217 Introduction 218 Experimental 219 Bulk Hyperbranch Incorporation 219 HB Polymers as Adhesion Promoters 220 Results and Discussion 221 Conclusion 227 References 228

217

Rheology of Hot-Melt PSAs: Influence of Polymer Structure 229 C. Derail and G. Marin Abstract 229 Introduction 229 Main Features of the Viscoelastic Behavior of the Pure Components, Blends, and Full Adhesive Formulations 231 Rheological Experiments 231 Rheological Behavior of the Pure Components: [SI], [SIS], and Pure Blends 231 Rheological Behavior of the Full Adhesive Formulations 233 A Model of the Rheological Behavior 236 A Model for the Pure Copolymers 236 A Model for the Blends [SIS–SI] 239 A Model for the Full Adhesive Formulations [SIS–SI–Resin] 239 Discussion 240

XIII

XIV

Contents

16.4.1 16.4.2 16.5

Molecular Design 240 On the Variation of the Secondary Elastic Plateau Modulus 241 Conclusions 245 Acknowledgments 247 References 247

17

Preparation and Characterization of UV-Crosslinkable Pressure-Sensitive Adhesives 249 H.-S. Do, S.-E. Kim, and H.-J. Kim Abstract 249 Introduction 249 Materials and Methods 252 Preparation of UV-Crosslinkable Acrylic PSA 252 Preparation of PSA Samples and UV Curing 253 FTIR-ATR Spectroscopy 253 DSC Measurement 254 PSA Performance 254 Results and Discussion 254 FTIR-ATR Measurements 254 PSA Performance 258 Probe Tack 258 Peel Strength 260 Shear Adhesion Failure Temperature (SAFT) 261 Conclusions 263 References 263

17.1 17.2 17.2.1 17.2.2 17.2.3 17.2.4 17.2.5 17.3 17.3.1 17.3.2 17.3.2.1 17.3.2.2 17.3.2.3 17.4

18

18.1 18.1.1 18.1.2 18.1.3 18.1.4 18.2 18.2.1 18.2.2 18.2.3 18.2.4 18.2.5 18.2.6

Contribution of Chemical Interactions to the Adhesion Between Evaporated Metals and Functional Groups of Different Types at Polymer Surfaces 265 J. Friedrich, R. Mix, and G. Kühn Abstract 265 Introduction 266 Interactions Between Metal Atoms and Functional Groups at Polymer Surfaces 266 Preparation of the Plasma-Modified Polymer Surfaces 267 Interactions Between Evaporated Al and Functional Groups 269 Adhesive Bond Strength and Concentration of Functional Groups 269 Materials and Methods 270 Materials 270 Plasma Pretreatment of Polymers 271 Deposition of Adhesion-Promoting Plasma Polymer Layers 271 Surface Analysis 271 Labeling of Functional Groups 272 Contact Angle Measurements 272

Contents

18.2.7 18.2.8 18.3 18.3.1 18.3.2 18.3.3 18.3.4 18.3.5 18.3.6 18.3.7 18.3.7.1 18.3.7.2 18.4 18.4.1 18.4.2 18.5

19

19.1 19.2 19.3 19.3.1 19.3.2 19.3.3 19.3.3.1 19.3.3.2 19.3.3.3 19.4

20

20.1 20.1.1

Metal Deposition 272 Peel Strength Measurements 273 Results 273 Production of Polymer Surfaces Containing Functional Groups 273 Surface Free Energy Measurements 275 Peel Strength Measurements of Al-Plasma Modified PP Systems 276 Peel Strength of Al–Plasma-Produced Homopolymer–PP Systems 277 Peel Strength of Al–Plasma Copolymer–PP Systems 277 Plasma Pretreatment of PTFE Surfaces 279 Peel Strength Measurements of Al–PTFE Systems 281 Hydrogen Plasma Pretreatment of PTFE 281 Hydrogen Plasma Pretreatment of PTFE and Deposition of Plasma Polymer Layers 281 Discussion 282 Contribution of Chemical Bonds to the Resulting Adhesion Strength 282 Dependence of Adhesion Strength on Concentration of Functional Groups at the Polymer S 284 Conclusion 285 References 286 Alkene Pulsed Plasma Functionalized Surfaces: An Interfacial Diels-Alder Reaction Study 289 F. Siffer, J. Schultz, and V. Roucoules Abstract 289 Introduction 289 Materials and Methods 290 Results and Discussion 292 Interfacial Chemistry 292 Cycloaddition 294 Kinetics 295 Monolayers 295 Plasma Polymer Thin Films 298 Comparison of Surface Reaction in Monolayers and Plasma Polymer Thin Films 299 Conclusion 302 References 303 Laser Surface Treatment of Composite Materials to Enhance Adhesion Properties 305 Q. Bénard, M. Fois, M. Grisel, and P. Laurens Introduction 305 Why Treat a Composite Surface? 305

XV

XVI

Contents

20.1.2 20.2 20.2.1 20.2.2 20.2.3 20.3 20.3.1 20.3.2 20.3.2.1 20.3.2.2 20.4

Available Treatments for Composite Surfaces 305 Materials and Methods 307 Composite Materials 307 Surface Analyses 307 Single Lap Shear Tests 308 Results and Discussion 308 Why Excimer Laser Treatment? 308 Excimer Laser Surface Treatment 310 Surface Characterization 310 Mechanical Tests 312 Conclusion 317 References 318

21

Effects of the Interphase on the Mechanical Behavior of Thin Adhesive Films – a Modeling Approach 319 S. Diebels, H. Steeb, and W. Possart Abstract 319 Introduction 319 Theoretical Framework 322 Applications and Examples 325 Uniaxial Tension Test 326 Simple Shear Test 330 Conclusion 330 References 333

21.1 21.2 21.3 21.3.1 21.3.2 21.4

22

22.1 22.2 22.3 22.4 22.5 22.6

23

23.1 23.2 23.2.1

Effect of the Diblock Content on the Adhesive and Deformation Properties of PSAs Based on Styrenic Block Copolymers 337 C. Creton, A. Roos, and A. Chiche Introduction 337 Block Copolymer Based Adhesives 339 Effect of the Diblock Content on Adhesive and Deformation Properties 348 Understanding the Structure of the Extended Foam 350 Interfacial Fracture 356 Summary 360 Acknowledgments 361 References 361 Contact Mechanics and Interfacial Fatigue Studies Between Thin Semicrystalline and Glassy Polymer Films R. L. McSwain, A. R. Markowitz, and K. R. Shull Abstract 365 Introduction 365 Materials and Methods 369 Materials and Sample Preparation 369

365

Contents

23.2.2 23.2.3 23.3 23.4 23.4.1 23.4.2 23.4.3 23.5

Pull-Off Test 371 Cyclic Interfacial Fatigue Test 374 Results 374 Discussion 381 Wetting Behavior and PEO/TMPC Miscibility 381 PEO/TMPC Interfacial Width and Adhesion 382 PDMS Rupture 384 Conclusion 385 Acknowledgments 385 References 385

24

Local and Global Aspects of Adhesion Phenomena in Soft Polymers 387 M.-F. Vallat Abstract 387 Introduction 387 The Molecular Interphase 388 Autohesion of Polyisoprene 389 Autoadhesion of EPDM 393 Macroscopic Interphases 395 Vulcanized Elastomers 395 Polyurethane Joints 398 Conclusion 400 References 401

24.1 24.2 24.2.1 24.2.2 24.3 24.3.1 24.3.2 24.4

25

25.1 25.2 25.3 25.4 25.5

26

26.1 26.2 26.3 26.4

Calibration and Evaluation of Nonlinear Ultrasonic Transmission Measurements of Thin-Bonded Interfaces 403 S. Hirsekorn, A. Koka, S. Kurzenhäuser, and W. Arnold Abstract 403 Introduction 403 Experimental and Calibration Procedure 404 Calibrated Ultrasonic Transmission Measurements 406 Ultrasonic Measurement and Destructive Tests 410 Conclusion 418 Acknowledgments 418 References 419 Debonding of Pressure-Sensitive Adhesives: A Combined Tack and Ultra-Small Angle X-Ray Scattering Study 421 E. Maurer, S. Loi, and P. Müller-Buschbaum Abstract 421 Introduction 421 In-Situ Small Angle Scattering Using Synchrotron Radiation 423 Microscopically Inaccessible Substructures 426 Conclusion 432

XVII

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Contents

Acknowledgments 433 References 433 27

27.1 27.2 27.3 27.4 27.5

28

28.1 28.2 28.2.1 28.2.2 28.3 28.3.1 28.3.2 28.3.3 28.4

29

29.1 29.2 29.3 29.3.1 29.3.2 29.3.3 29.4

Nondestructive Testing of Adhesive Curing in Glass–Metal Compounds by Unilateral NMR 435 K. Kremer, B. Blümich, F.-P. Schmitz, and J. Seitzer Abstract 435 Introduction 436 Nuclear Magnetic Resonance (NMR) and the NMR-MOUSE 436 Quality Control 437 Application 438 Conclusion and Outlook 442 Acknowledgments 442 References 443 Chemical Processes During Aging in Ultra-thin Epoxy Films on Metals 445 A. Meiser, C. Wehlack, and W. Possart Abstract 445 Introduction 445 Experimental 447 Sample Preparation 447 Aging Conditions 447 Results and Discussion 448 Crosslinking 448 Additional Aging Effects 451 Band Assignment and Chemical Aging Processes 458 Conclusion 462 Acknowledgments 463 References 463 Depth-Resolved Analysis of the Aging Behavior of Epoxy Thin Films by Positron Spectroscopy 465 J. Kanzow, F. Faupel, W. Egger, P. Sperr, G. Kögel, C. Wehlack, A. Meiser, and W. Possart Abstract 465 Introduction 465 Materials and Methods 466 Results 467 PALS Investigation of an Unaged Epoxy Film 468 PALS Investigation of Aged Epoxy Films 469 Further Investigations of Aged Epoxy Films 471 Discussion and Conclusion 474 Acknowledgments 476 References 476

Contents

30

30.1 30.2 30.2.1 30.2.2 30.2.3 30.2.4 30.3 30.3.1 30.3.1.1 30.3.1.2 30.3.1.3 30.3.2 30.3.2.1 30.3.2.2 30.3.2.3 30.4

31

31.1 31.2 31.3 31.3.1 31.3.2 31.3.2.1 31.3.2.2 31.3.3 31.3.4

Epoxies on Stainless Steel – Curing and Aging 479 D. Fata, C. Bockenheimer, and W. Possart Abstract 479 Introduction 480 Materials and Methods 481 Materials 481 Sample Preparation 482 Aging Experiments 482 Characterization of Aged Specimens 483 Results and Discussion 484 The RT Curing Epoxy System (EP1) 484 Curing of EP1 484 Thermal Aging of EP1 after Post-Curing at 408C 487 Hydro-thermal Aging of EP1 492 The Hot-Curing Epoxy System (EP2) 495 Curing of EP2 495 Thermal Aging of EP2 498 Hydro-thermal Aging of EP2 500 Conclusion 503 Acknowledgment 505 References 505 Scanning Kelvin Probe Studies of Ion Transport and De-adhesion Processes at Polymer/Metal Interfaces 507 K. Wapner and G. Grundmeier Abstract 507 Introduction 508 Theory and Experimental Set-Up of a Scanning Kelvin Probe 509 Applications of Scanning Kelvin Probe Studies in Adhesion Science 514 Diffusion of Ions into Metal/Adhesive Interfaces 514 Corrosive Degradation of the Polymer/Metal Interface 516 Cathodic Delamination on Adhesive-Coated Iron 516 Anodic Delamination (Filiform Corrosion) on Coated Aluminum 518 Detection of Wet Debonding 520 A New Scanning Kelvin Probe Blister Test 521 Acknowledgment 523 References 523

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32

32.1 32.2 32.2.1 32.2.2 32.2.3 32.3 32.3.1 32.3.2 32.4 32.4.1 32.4.2 32.4.3 32.5 32.5.1 32.5.2 32.6

33

33.1 33.2 33.3 33.4 33.4.1 33.4.2 33.4.3 33.4.4 33.5 33.6

Advanced Mass Transport Applications with Elastic Bonding of Sandwich Components 525 S. Koch, A. Starlinger, and X. Wang Abstract 525 Introduction 525 Stress Distribution in Different Joints 526 Stress Distribution in Bolted Joints 527 Stress Distribution in a Stiff Adhesive Joint 528 Stress Distribution in an Elastic Adhesive Joint 529 Applications of Flexible Adhesives in Mass Transportation Systems 529 GRP Front Cab 530 Application in Tram Design 530 Methods of Modeling Flexible Adhesives 531 Modeling Methods for Detailed Local Analysis 532 Modeling Methods for Large Global Structural Analysis 533 Comparison of the TR08 Results from FE Analysis and from Measurement on Lathen Test Track 534 Joint Design, Production, and Testing 535 Production of Adhesive Joints 536 Joint Testing 536 Conclusion 537 References 537 Adhesive Joints for Modular Components in Railway Applications 539 C. Nagel, M. Brede, M. Calomfirescu, J. Sauer, E. A. Ullrich, T. Fertig, and O.-D. Hennemann Abstract 539 Introduction 539 Adhesives and Adherends 540 Surface Pretreatment 541 Mechanical Behavior of Adhesives and Joints 542 Elastic–Plastic Properties of Structural Adhesive Systems 543 Hyperelastic Properties of Flexible Adhesive Systems 544 Creep Behavior of Adhesive Joints 545 Fatigue Properties of Adhesive Joints 547 Environmental Influences and Design of Structures 550 Conclusion 553 Acknowledgment 553 References 554

Contents

34

34.1 34.2 34.2.1 34.2.2 34.2.3 34.2.4 34.2.5 34.3 34.3.1 34.3.2 34.3.3 34.3.4 34.3.5 34.5

Behavior of Dismantlable Adhesives Including Thermally Expansive Microcapsules 555 Y. Nishiyama and C. Sato Abstract 555 Introduction 555 Materials and Methods 557 Materials 557 Volume Expansion of the Cured Bulk Adhesive 558 Dismantlability of Joints Bonded with the Dismantlable Adhesive 559 Bond Strength of the Dismantlable Adhesive 559 PVT (Pressure–Volume–Temperature) Tests 560 Results and Discussion 561 Volume Expansion of the Cured Bulk Adhesive 561 Dismantlability of Joints Bonded with the Dismantlable Adhesive 562 Bond Strength of the Dismantlable Adhesive 564 PVT Relationship of Microcapsules and Dismantlable Adhesive 565 Discussion 567 Conclusion 567 References 568 Subject Index

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List of Contributors M. Amkreutz Universität Paderborn Fachbereich Physik Theoretische Physik Warburger Str. 100 33100 Paderborn Germany

U. Bakowsky Department of Pharmaceutical Technology and Biopharmaceutics Philipps University Marburg Postfach 35032 Marburg Germany

W. Arnold Fraunhofer-Institut für Zerstörungsfreie Prüfverfahren (IZFP) 66123 Saarbrücken Germany

J. Baller Laboratoire Européen de Recherche Universitaire Saarland-Lorraine (LERUSL) Université du Luxembourg Laboratoire de Physique des Matériaux 162a, avenue de al Faiencerie 1511 Luxembourg

M. Aufray INSA de Lyon Laboratoire des Matériaux Macromoléculaires (IMP/LMM) 20, av Albert Einstein 69621 Villeurbanne Cedex France R. Bactavatchalou Laboratoire Européen de Recherche Universitaire Saarland-Lorraine (LERUSL) Universität des Saarlandes Fakultät für Physik und Elektrotechnik 7.2 Gebäude 38 Postfach 151150 66041 Saarbrücken Germany

Q. Bénard Unité de Recherche en Chimie Organique et Macromoléculaire EA3221 Université du Havre 25, rue Philippe-Lebon 76600 Le Havre France S. Bistac Institut de Chimie des Surfaces et Interfaces CNRS-ICSI 15, rue Jean Starcky 68057 Mulhouse Cedex France


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List of Contributors

A. T. Blumenau Max-Planck-Institut für Eisenforschung Max-Planck-Str. 1 40237 Düsseldorf Germany B. Blümich ITMC RWTH Aachen Worringer Weg 1 52074 Aachen Germany C. Bockenheimer Universität des Saarlandes Gebäude 22-6 Postfach 151150 66041 Saarbrücken Germany T. Brand Rafael P.O. Box 2250 Haifa 31021 Israel M. Brede Fraunhofer-Institut für Fertigungstechnik und Angewandte Materialforschung (IFAM) Wiener Str. 12 28359 Bremen Germany M. Brogly Université de Haute Alsace (UHA) and Institut de Chimie des Surfaces et Interfaces (ICSI) CNRS UPR 9069 15, rue Jean Starcky 68057 Mulhouse Cedex France

A. Buchman Rafael P.O. Box 2250 Haifa 31021 Israel M. Calomfirescu Siemens AG Transportation Systems Duisburger Str. 145 47829 Krefeld Germany G. Castelein Université de Haute Alsace (UHA) and Institut de Chimie des Surfaces et Interfaces (ICSI) CNRS UPR 9069 15, rue Jean Starcky 68057 Mulhouse Cedex France A. Chiche Laboratoire de Physico-Chimie des Polymères et Milieux Dispersés UMR 7615 ESPCI 10, rue Vauquelin 75231 Paris Cédex 05 France J. Chung Bundesanstalt für Materialforschung und -prüfung (BAM) Div. VI.2 Unter den Eichen 87 12205 Berlin Germany C. Creton Laboratoire de Physico-Chimie des Polymères et Milieux Dispersés UMR 7615 ESPCI 10, rue Vauquelin 75231 Paris Cédex 05 France


C. Derail Laboratoire de Physico-Chimie des Polymères UMR-CNRS 5067 Université de Pau et des Pays de l’Adour – CURS Avanue de l’Université 64013 Pau France S. Diebels Lehrstuhl für Technische Mechanik Universität des Saarlandes Gebäude 22-12 Postfach 151150 66041 Saarbrücken Germany H.-S. Do Adhesion & Bio-Composites Laboratory Program in Environmental Materials Science Seoul National University Seoul 151-921 Republik of Korea H. Dodiuk-Kenig Shenkar College of Engineering & Design Ramat-Gan Israel W. Egger Universität der Bundeswehr München Institut für Nukleare Festkörperphysik Werner-Heisenberg-Weg 85577 Neubiberg Germany T. Elzein Institut de Chimie des Surfaces et Interfaces 15, rue Jean Starcky 68057 Mulhouse France

C. Ehrhardt Department of Biopharmaceutics and Pharmaceutical Technology Saarland University 66123 Saarbrücken Germany D. Fata Universität des Saarlandes Gebäude 22-6 Postfach 151150 66041 Saarbrücken Germany F. Faupel Lehrstuhl für Materialverbunde Technische Fakultät Christian-Albrechts-Universität Kiel Kaiserstr. 2 24143 Kiel Germany T. Fertig Henkel Teroson GmbH Henkel-Teroson-Str. 57 69123 Heidelberg Germany M. Fois Unité de Recherche en Chimie Organique et Macromoléculaire EA3221 Université du Havre 25, rue Philippe-Lebon 76600 Le Havre France T. Frauenheim Universität Paderborn Fachbereich Physik Theoretische Physik Warburger Str. 100 33100 Paderborn Germany

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J. Friedrich Bundesanstalt für Materialforschung und -prüfung (BAM) Unter den Eichen 87 12205 Berlin Germany

O.-D. Hennemann Fraunhofer-Institut für Fertigungstechnik und Angewandte Materialfoschung (IFAM) Wiener Str. 12 28359 Bremen Germany

A. Galliano Institut de Chimie des Surfaces et Interfaces CNRS-ICSI 15, rue Jean Starcky 68057 Mulhouse Cedex France

S. Hirsekorn Fraunhofer-Institut für Zerstörungsfreie Prüfverfahren (IZFP) 66123 Saarbrücken Germany

Z. Gold Israel Plastics and Rubber Center Ramat-Gan Isreal

D. Hoekstra Department of Membrane Cell Biology University of Groningen 9700 AV Groningen Netherlands

M. Grisel Unité de Recherche en Chimie Organique et Macromoléculaire EA3221 Université de Havre 25, rue Philippe-Lebon 76600 Le Havre France G. Grundmeier Department for Interface Chemistry and Surface Engineering Max-Planck-Institut für Iron Research Max-Planck-Str. 1 40237 Düsseldorf Germany A. Hartwig Fraunhofer-Institut für Fertigungstechnik und Angewandte Materialfoschung (IFAM) Wiener Str. 12 28359 Bremen Germany

D. Jahn Institute of Microbiology Technical University of Braunschweig 38106 Braunschweig Germany G. Kalinka Bundesanstalt für Materialforschung und -prüfung (BAM) Div. VI.2 Unter den Eichen 87 12205 Berlin Germany J. Kanzow Lehrstuhl für Materialverbunde Technische Fakultät Christian-Albrechts-Universität Kiel Kaiserstr. 2 24143 Kiel Germany


S. Kenig Shenkar College of Engineering & Design Ramat-Gan Israel

A. Koka Fraunhofer-Institut für Zerstörungsfreie Prüfverfahren (IZFP) 66123 Saarbrücken Germany

H.-J. Kim Adhesion & Bio-Composites Laboratory Program in Environment Materials Science Seoul National University Seoul 151-921 Republic of Korea

K. Koschek Fraunhofer-Institut für Fertigungstechnik und Angewandte Materialfoschung (IFAM) Wiener Str. 12 28359 Bremen Germany

S.-E. Kim Adhesion & Bio-Composites Laboratory Program in Environmental Materials Science Seoul National University Seoul 151-921 Republic of Korea C. Kneuer Leipzig University Department of Pharmacology Pharmacy and Toxicology An den Tierkliniken 15 04103 Leipzig Germany S. Koch Sika Technology AG Tüffenwies 16 8048 Zurich Switzerland G. Kögel Universität der Bundeswehr München Institut für Nukleare Festkörperphysik Werner-Heisenberg-Weg 85577 Neubiberg Germany

K. Kremer ITMC RWTH Aachen Worringer Weg 1 52074 Aachen Germany J. K. Krüger Laboratoire Européen de Recherche Universitaire Saarland-Lorraine (LERUSL) Universität des Saarlandes Fakultät für Physik und Elektrotechnik 7.2 Gebäude 38 Postfach 151150 66041 Saarbrücken Germany T. Krüger Universität Paderborn Fachbereich Physik Theoretische Physik Warburger Str. 100 33100 Paderborn Germany

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G. Kühn Bundesanstalt für Materialforschung und -prüfung (BAM) Unter den Eichen 87 12205 Berlin Germany S. Kurzenhäuser Fraunhofer-Institut für Zerstörungsfreie Prüfverfahren (IZFP) 66123 Saarbrücken Germany P. Laurens Cooperative du Laser Franco-Allemand Laboratoire d’Application des Lasers de Puissance Direction Générale de l’Armement 94114 Arcueil Cedex France G. Legeay Centre de Transfert de Technologies du Mans rue Thalès de Milet 72000 Le Mans France C.-M. Lehr Department of Biopharmaceutics and Pharmaceutical Technology Saarland University 66123 Saarbrücken Germany P. Li Department of Biopharmaceutics and Pharmaceutical Technology Saarland University 66123 Saarbrücken Germany

D. Liebschner Laboratoire Européen de Recherche Universitaire Saarland-Lorraine (LERUSL) Universität des Saarlandes Fakultät für Physik und Elektrotechnik 7.2 Gebäude 38 Postfach 151150 66041 Saarbrücken Germany C. Loehbach Department of Biopharmaceutics and Pharmaceutical Technology Saarland University 66123 Saarbrücken Germany S. Loi TU München Physikdepartment E13 James-Franck-Str. 1 85748 Garching Germany A. Lühring Fraunhofer-Institut für Fertigungstechnik und Angewandte Materialfoschung (IFAM) Wiener Str. 12 28359 Bremen Germany G. Marin Laboratoire de Physico-Chimie des Polymères UMR-CNRS 5067 Université de Pau et des Pays d l’Adour – CURS Avenue de l’Université 64013 Pau France


R. Markowitz Department of Materials Science and Engineering Northwestern University 2220 Campus Dr. Evanston IL 60208-3108 USA E. Maurer TU München Physikdepartment E13 James-Franck-Str. 1 85748 Garching Germany R. L. McSwain Department of Materials and Engineering Northwestern University 2220 Campus Dr. Evanston IL 60208-3108 USA A. Meiser Universität des Saarlandes Gebäude 22-6 Postfach 151150 66041 Saarbrücken Germany R. Mix Bundesanstalt für Materialforschung und -prüfung (BAM) Unter den Eichen 87 12205 Berlin Germany R. Mühlhaupt Freiburger Materialforschungszentrum (FMF) und Institut für Makromolekulare Chemie of the Albert Ludwigs University Freiburg Stefan-Meier-Straße 31 79104 Freiburg Germany

U. Müller Laboratoire Européen de Recherche Universitaire Saarland-Lorraine (LERUSL) Universität des Saarlandes Fakultät für Physik und Elektrotechnik 7.2 Gebäude 38 Postfach 151150 66041 Saarbrücken Germany P. Müller-Buschbaum TU München Physikdepartment E13 James-Franck-Str. 1 85748 Garching Germany M. Munz Bundesanstalt für Materialforschung und -prüfung (BAM) Div. VI.2 Unter den Eichen 87 12205 Berlin Germany C. Nagel Fraunhofer-Institut für Fertigungstechnik und Angewandte Materialfoschung (IFAM) Wiener Str. 12 28359 Bremen Germany Y. Nishiyama Interdisciplinary Graduate School of Science and Engineering Tokyo Institute of Technology 4259 Nagatsuta Midori-ku Yokohama 226-8503 Japan

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O. Noel Université de Haute Alsace (UHA) Institut de Chimie des Surfaces et Interfaces (ICSI) CNRS UPR 9069 15, rue Jean Starcky 68057 Mulhouse Cedex France

V. Roucoules Institut de Chimie des Surfaces et Interfaces ICSI-CNRS-UPR 9069 15, rue Jean Starcky BP 2488 68057 Mulhouse Cedex France

F. Poncin-Epaillard Laboratoire Polymères Colloïdes et Interfaces – UMR CNRS 6120 Université du Maine Avenue O Messiaen 72085 Le Mans France

D. Rouxel Laboratoire Européen de Recherche Universitaire Saarland-Lorraine (LERUSL) Université Henri Poincaré, Nancy 1 Boulevard des Aiguillettes 54506 Vandoeuvre les Nancy France

W. Possart Laboratoire Européen de Recherche Universitaire Saarland-Lorraine (LERUSL) Universität des Saarlandes Gebäude 22-6 Postfach 151150 66041 Saarbrücken Germany A. A. Roche INSA de Lyon Laboratoire des Matériaux Macromoléculaires (IMP/LMM) 17, rue Jean Capelle 69621 Villeurbanne Cedex France A. Roos Laboratoire de Physico-Chimie des Polymères et Milieux Dispersés UMR 7615 ESPCI 10, rue Vauquelin 75231 Paris Cédex 05 France

M. Sander Laboratoire Européen de Recherche Universitaire Saarland-Lorraine (LERUSL) Universität des Saarlandes Fakultät für Physik und Elektrotechnik 7.2 Gebäude 38 Postfach 151150 66041 Saarbrücken Germany C. Sato Precision and Intelligence Laboratory Tokyo Institute of Technology 4259 Nagatsuta Midori-ku Yokohama 226-8503 Japan J. Sauer Huntsman Advanced Materials (Switzerland) GmbH Klybeckstr. 200 4057 Basel Switzerland


P. Schiffels Fraunhofer-Institut für Fertigungstechnik und Angewandte Materialfoschung (IFAM) Wiener Str. 12 28359 Bremen Germany F.-P. Schmitz Poly-BEADS Am Butterberg 20 21335 Lüneburg Germany B. Schneider Fraunhofer-Institut für Fertigungstechnik und Angewandte Materialfoschung (IFAM) Wiener Str. 12 28359 Bremen Germany J. Schultz Université de Haute Alsace (UHA) and Institut de Chimie des Surfaces et Interfaces (ICSI) CNRS UPR 9069 15, rue Jean Starcky 68057 Mulhouse Cedex France J. Seitzer EFTEC AG Hofstr. 31 8590 Romanshorn Switzerland K. R. Shull Department of Materials Science and Engineering Northwestern University 2220 Campus Dr. Evanston IL 60208-3108 USA

F. Siffer Institut de Chimie des Surfaces et Interfaces ICSI-CNRS-UPR 9069 15, rue Jean Starcky BP 2488 68057 Mulhouse Cedex France P. Sperr Universität der Bundeswehr München Institut für Nukleare Festkörperphysik Werner-Heisenberg-Weg 85577 Neubiberg Germany A. Starlinger Stadler Altenrhein AG Park Altenrhein für Industrie und Gewerbe 9423 Altenrhein Switzerland H. Steeb Lehrstuhl für Technische Mechanik Universität des Saarlandes Gebäude 22-12 Postfach 151150 66041 Saarbrücken Germany E. A. Ullrich Henkel Teroson GmbH Henkel-Teroson-Str. 75 69123 Heidelberg Germany M.-F. Vallat Institut de Chimie des Surfaces et Interfaces – UPR 9069 CNRS et Université de Haute-Alsace 15, rue Jean Starcky BP 2488 68057 Mulhouse Cedex France

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X. Wang Alcan Alesa Engineering Ltd. Alcan Mass Transportation SystemP.O. Box 1250 8048 Zurich Switzerland

C. Wehlack Universität des Saarlandes Gebäude 22-6 Postfach 151150 66041 Saarbrücken Germany

K. Wapner Department of Interface Chemistry and Surface Engineering Max-Planck-Institute for Iron Research Max-Planck-Str. 1 40237 Düsseldorf Germany

M. Wiegemann Fraunhofer Institute for Manufacturing Engineering and Applied Materials (IFAM) Bonding Technology and Surfaces Department Wiener Str. 12 28359 Bremen Germany

J. F. Watts The Surface Analysis Laboratory School of Engineering University of Surrey Guildford Surrey GU2 7XH UK

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1 The Interfacial Chemistry of Adhesion: Novel Routes to the Holy Grail? J. F. Watts

Abstract

This paper reviews the progress that has been made in the use of surface analytical techniques such as XPS and ToF-SIMS to obtain chemical information from the buried interface of an organic coating or adhesives system. Such a task is nontrivial, as the interfacial region, at most a few nanometers wide, is buried between many tens or hundreds of micrometers of substrate and overlayer. Two methodologies are described for the determination of interface chemistry: the deposition of a very thin (ca. 2 nm) film of coating or adhesive on the substrate followed by the use of XPS or ToF-SIMS to “look” through the layer to extract chemical information specific to the substrate; and the use of chemical or mechanical sectioning to present the sample in such a way that the interface region of the system is submitted for analysis. In the former category, examples are given of the deposition of an organosilane on aluminum, which yields a covalent Al–O–Si bond, and the subsequent extension of covalency via a TDI urone (toluene diisocyante urone) curing agent and an epoxy resin into the bulk of the cured adhesive. The continuity of covalent bonding from the substrate to the bulk resin ensures good durability of such a system. Oxide stripping and ultra-low-angle microtomy techniques are used to provide examples of the development of unique interface chemistry and the elucidation of interface concentration gradients on a duplex organic coatings system. It is concluded that such approaches have much to offer the adhesion scientist in the search for the Holy Grail: the ability to reverse-engineer interface chemistry in order to confer specific properties.

1.1 Introduction

In any attempt to understand the physico-chemical phenomena responsible for adhesion between a polymer phase, such as an adhesive or organic coating, and a solid substrate, there is a need to be able to access the interface or interphase Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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1 The Interfacial Chemistry of Adhesion: Novel Routes to the Holy Grail?

Fig. 1.1 The problem that exists in the analysis of the buried interface: a region responsible for adhesion that is nanometers thick buried between thick layers of adhesive or coating and substrate materials.

between the two materials. This is easier said than done; the complexity of the situation is indicated in Fig. 1.1, in that significant thicknesses (many tens of micrometers at the very least) of both substrate and overlayer obscure the adhesion layer or boundary that exists between the two. In order to examine this critical region one must resort to the removal of large amounts of material by mechanical or chemical means, or construct model specimens in which the interphase region is more readily accessible by analytical methods such as XPS [1, 2] and ToF-SIMS [3]. The simple expedient of forensic analysis of a failure interface produced by mechanical, chemical, or electrochemical means is unlikely to provide a specimen that will yield the required information. Although it may lead to a greater understanding of the mechanism of failure, fundamental aspects relating to adhesion phenomena will invariably be elusive. The nature of interfacial bonding will, to a certain extent and in a very simplistic manner, influence joint strength. At this level the contributions to interfacial bonding are quite simply the number of bonds in place per unit area and the strength (more strictly the bond energy) of the various bond types. Thus it is possible to envisage the equivalence of the contribution to bond strength from many weak bonds (or rather less strong bonds) and fewer, stronger bonds, as indicated in Fig. 1.2 (a). This is all well and good, but when one starts to consider the resistance of the interface to aqueous exposure it is often found that strong bonds (such as covalent bonds) are much more resistant to hydrodynamic displacement than weaker bonds such as those of the van der Waals type (Fig. 1.2 b). Thus both the aeric density of bonding sites and a knowledge of bond type are essential if one is to understand the interfacial chemistry and its effect on adhesion and – more importantly – durability. In the longer term the aim is to be able to engineer specific chemistry at the interface which will provide the required level of performance from an adhesive joint or organic coating. This paper will build on previous reviews which have sought to explore the manner in which surface analysis methods can be purposefully employed to understand adhesion phenomena [4–6], with an emphasis on the elucidation of interphase chemistry. The rationale behind such an approach is that it is this critical region of a polymer/metal or polymer/polymer couple that will influence the performance of the overall system, be it the durability of an adhesive joint or the corrosion protection performance of an organic coating.

1.2 Development of a Model Interphase

Fig. 1.2 (a) Schematic view of the equivalence, in terms of adhesion, of many weak bonds and a few strong bonds. (b) The stronger (covalent) bonds are much more resistant to displacement by water than the weaker intermolecular bonds.

In essence there are two potential ways in which the interphase region can be approached; either by the use of systems based on real adhesives or organic coatings to create a model interphase, or by the sectioning, by some means or other, to expose the interphase region prior to analysis by XPS or ToF-SIMS. In this paper the use of both approaches, which have been widely explored in the author’s laboratory over the last three decades, will be described.


This route to interphase analysis may also be described as the thin-film approach, inasmuch as the basic principle involves the deposition of extremely thin layers (< 2 nm) of the mobile phase onto a substrate. In the work described here the mobile phase is the organic component of the system, but the same approach is applicable to studies of the metallization of polymer substrates. It is also convenient to combine studies of polymer interactions with solid substrates with studies of the adsorption characteristics of the organic components themselves. Such an approach has much to offer in adhesion research and the basis of studies of adsorption from a liquid phase and its applicability in adhesion has been discussed in detail elsewhere [7] so it will not be treated in depth here. A brief overview will, however, provide a background to this approach. The determination of gas-phase adsorption isotherms is a well-known methodology in surface chemistry; in this manner it is possible to describe adsorption as following Langmuir or other characteristic adsorption types. The conventional method of studying the adsorption of molecules from the liquid phase is to establish the depletion of the adsorbate molecule from the liquid phase. However, as first pointed out by Castle and Bailey [8], with the advent of surface analysis methods it is now

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very straightforward to monitor the actual uptake of the adsorbate on the solid surface by XPS or SIMS. The experiment itself is quite simple in that a set of coupons of the solid substrates are exposed to a series of dilute solutions of the candidate adsorbate. The uptake curve (the adsorption isotherm) is quite simply a plot of surface concentration of the adsorbate (determined by XPS or SIMS) versus the solution concentration. In the case of chemisorption of the adsorbate on the substrate the uptake curve will quickly reach a plateau, indicating that all the potential adsorption sites on the substrate are occupied by adsorbate molecules. An example of this is shown in Fig. 1.3, which indicates the adsorption of the diglycidyl ether of bisphenol A (DGEBA) on silane-treated aluminum. The measure of the uptake of the DGEBA is taken as the relative peak intensity of the mass 135 peak in the positive ToF-SIMS spectrum, which is very characteristic of the DGEBA molecule [9]. If necessary the thickness of the overlayer at monolayer coverage (the plateau region) can be determined by XPS; similarly, the adsorption regime which best describes the adsorption characteristics (that is, whether adsorption conforms to the Langmuir, Temkin, or another type) may be determined by simple diagnostic tests, as described elsewhere [7]. Although there are several possibilities, experience has shown that many of the adsorption phenomena of importance in adhesion are characterized by Langmuir adsorption, indicating an equivalence of adsorption sites (in terms of enthalpy of adsorption) on the solid substrate. Although the adsorption isotherm provides us with much important evidence regarding the aeric density of bonding sites and the type of adsorption that occurs, it tells us little about the interfacial reactions responsible for adhesion. In order to achieve this goal it is necessary, as indicated earlier, to examine a thin layer of the adsorbate on the substrate. The choice of such a specimen can be made from an adsorption isotherm if chemisorption is known to occur (indicated by a curve of the form of Fig. 1.3). A specimen taken from the plateau region of such an uptake

Fig. 1.3 Adsorption isotherm of DGEBA on GPS-treated aluminum derived from ToF-SIMS data. RPI is the relative peak intensity of the SIMS fragment of interest. This is equivalent to the normalised peak intensity (courtesy of Dr. A. Rattana). (GPS: c-glycidoxypropyltrimethoxysilane.)


curve will yield the maximum number of interfacial bonds with the minimum overlayer thickness. It should thus, in principle, be possible to use XPS or SIMS to probe the interfacial chemistry directly. Although this is perhaps the optimum approach, it is sometimes not possible to obtain the entire uptake curve and in such cases a very dilute solution should be used. An interesting study using the thin-film approach is provided by the adsorption of poly(methyl methacrylate) on a series of oxidized metal substrates [10]. By careful examination of the XPS C1s spectra it was possible to relate small changes in the relative position of the methoxy and ester components to specific interactions between the polymer and the metal substrates. The type of bonding observed depends strongly on the acido–basic properties of the metal oxide. The adsorption isotherms from these systems were not simple to interpret, as polymer conformation changes as the solution concentration increased gave the erroneous appearance of multilayer deposition [11]. An elegant example of this type of approach is the recent study of the interaction of an organosilane adhesion promoter (c-glycidoxypropyltrimethoxysilane, GPS) on aluminum. The concept of a formal covalent bond between aluminum and the organosilane is not new and was first suggested back in the 1970s, but it is only unambiguously identifiable using high-resolution ToF-SIMS. The spectrum of Fig. 1.4 is a high-resolution mass spectrum of nominal mass m/z = 71. The intense peak labelled SiOAl+ is indicative of the bonding scheme shown in Scheme 1.1 [12]. It should be noted that, in this case, the organosilane was applied to the metallic substrate as a primer (1% aqueous solution) and then cured at 93 8C for 30 min. Evidence is emerging from current work that when the organosilane is included in the formulation of a room temperature curing adhesive, such a reaction is not present at the interface [13]. It seems that there is a subtle synergy between curing agent and organosilane, leading to an interpenetrating network. In a similar vein ToF-SIMS has also been used identify the interaction between an aminosilane and iron surfaces [14], but this interaction does seem to occur at ambient temperature. In dealing with commercial systems the investigative scientist is faced with many potential problems but the most significant is that such systems will, in the main, be very complex formulations of many individual components, and for obvious reasons associated with commercial sensitivity, it is unlikely that those outside the manufacturing company will be privy to details regarding the formulation. A typical structural adhesive will have many components in the formulation, some of which are [15]: · liquid epoxy resin · solid epoxy resin · polymeric modifiers · hardener · accelerator · fillers and additives · pigments and dyestuffs · support carrier.

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Fig. 1.4 High-resolution ToF-SIMS spectrum of the m/ = 71 region showing the intense component assigned to SiOAl+, indicative of the formation of a covalent bond (courtesy of Dr. M.-L. Abel).

Given this complexity, there is really no option but to investigate some components in isolation before going on to the fully formulated product. In work on an aerospace structural adhesive, the adsorption characteristics of the liquid epoxy resin (DGEBA) and the curing agent (toluene diisocyante urone, TDI urone) were studied independently on bare aluminum and aluminum treated with GPS. Then the interactions of dilute solutions of the adhesive with the two substrates were studied, enabling a detailed model of the interfacial chemistry to be proposed. The adsorption isotherms for the TDI urone curing agent adsorbed from a dilute solution of the adhesive are presented in Fig. 1.5. It is clear that the adsorption on the bare aluminum surface appears to be twice that of the GPS-treated aluminum. If the intensity of the halved curve obtained from the bare aluminum is halved it is found that it is coincident with that

1.2 Development of a Model Interphase Scheme 1.1 Bonding between the hydrated aluminum substrate and GPS responsible for the formation of the type of bond seen in the ToF-SIMS spectrum of Fig. 1.4.

from the GPS-treated substrate. In order to resolve this apparent conundrum it is helpful to consider the source of the ion used to provide the data of Fig. 1.5. The structure of the curing agent is shown in Scheme 1.2 and the ion at m/ z = 58 is assigned to the CH3–NH–C = O+ ion that is readily generated at either end of the molecule. The reaction scheme that is thought to occur (Scheme 1.3) shows the manner in which a GPS molecule, bonded to an aluminum substrate, might interact with a curing agent molecule. The immobilization of the TDI urone by reaction with the oxirane ring of the GPS molecule will mean that it is less likely that the bonded end of the amine will yield the characteristic m/z = 58 fragment. In the case of the bare aluminum the curing agent will interact by way of acid–base interactions via, for example, the carbonyl group of the curing agent molecule; as the interaction is not so strong as the covalent

Fig. 1.5 Uptake of the TDI urone curing agent from a dilute solution of the commercial adhesive of a bare (grit-blasted) aluminum substrate and a similar substrate treated with GPS. When the curve obtained from the bare aluminum (no GPS) substrate is divided by two it is exactly coincident with the data from the GPS-treated substrate (courtesy of Dr. A. Rattana).

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bond described above, both ends of the molecule are available to yield the characteristic ion. Further evidence for such an interaction with the GPS-treated substrate is provided by way an ion at m/z = 277 which is formed by scission adjacent to the CH–OH group formed on opening of the oxirane ring [16]. This type of investigation provides important information regarding the type of bonding that forms at the interface and can be represented in schematic form: Scheme 1.4 indicates a formal covalent bonding from substrate, through adhesion promoter, curing agent, DGEBA, into the bulk of the cross-linked adhesive. Given the comments regarding the hydrodynamic stability of covalent bonds relative to secondary bonds exemplified by the schematic of Fig. 1.2, one would expect superior performance from such a system, and this is indeed the case.

Scheme 1.2 Fragmentation of the TDI urone molecule to yield the m/z = 58 fragment in the positive ToF-SIMS spectrum.

1.3 The Buried Interface

Although the use of the thin-film method to provide a model interphase for analysis has much to commend it, the analysis of such a region formed between adhesive, coating, and substrate is perhaps more attractive in a number of situations. There are a number of options involving electron microscopy and surface analysis but in all cases the specimen preparation is the key to the optimum results. In the case of electron microscopy the most obvious expedient, the use a metallographic cross-section in conjunction with scanning electron microscopy (SEM) and energy-dispersive X-ray analysis (EDX) is ruled out for a number of reasons. Polishing a cross-section will invariably lead to smearing of


Scheme 1.3 Proposed interaction for the reaction of the TDI urone with a GPS-treated aluminum surface.

Scheme 1.4 The interaction mechanism of the adhesive with the GPS molecule.

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the polymer phase which will reduce the level of information attainable, and although electron microscopy can be readily carried out the interaction volume of the electron beam with the sample will mean that analytical resolution will be of the order of one micrometer, not nearly good enough to supply an analysis of the interfacial region relevant to adhesion. The solution to these problems is twofold; the smearing is minimized by cutting sections with a diamond knife, and the interaction volume is reduced by making the specimen thinner. In short, one carries out analytical transmission electron microscopy on specimens prepared by ultramicrotomy. Analysis can be carried out by EDX or parallel electron energy loss spectroscopy (PEELS) in the imaging mode, referred to as energy-filtered TEM (EFTEM). Although EDX provides only an elemental analysis, PEELS does (in theory) provide chemical state information analogous to that from XPS. Examples of the latter are few and far between, but the use of TEM for microscopy of the adhesive joints with a view to establishing interpenetration of adhesive with pretreatment layers is well established. Fig. 1.6 shows EFTEM data from adhesive-bonded anodised aluminum treated with and without a primer layer for a study of this type. It is clear that in the case of the substrate treated with the primer (Fig. 1.6 a) the organic phase (indicated by the carbon map) penetrates deep into the porous structure of the anodic layer, while in the absence of the primer (Fig. 1.6 b) complete penetration is not achieved [17]. The level of chemical information that can be gleaned from XPS and ToF-SIMS is potentially much greater than with analytical TEM but at the expense of spatial resolution. Rather than using a metallographic cross-section one must adopt a “plan view” specimen orientation, or something close to this geometry. One approach that has been used with a degree of success is to remove the metal sub-

Fig. 1.6 Energy-filtered (PEELS) TEM images of adhesively bonded aluminum: (a) the interpenetration of organic and oxide phases that is achieved when a primer is used; (b) in the absence of a primer, the adhesive merely forms an interfacial boundary with the oxide [17].


strate chemically, but not the oxide layer, and then mount the duplex polymer/oxide film for analysis in the spectrometer with the oxide side uppermost. By sequential sputter removal combined with analysis (sputter depth profiling) it is possible to produce a compositional depth profile toward the polymer/metal oxide interface, as indicated in Fig. 1.7. Once the interfacial region is reached there will be potential problems of degradation of the polymer phase but in many cases it is possible to extract information relating to the manner of oxide/polymer interaction. Important issues are the choice of stripping reagent: iodine in methanol works well for steels, as does NaOH for aluminum substrates. An example of this type of investigation, taken from the study of a polybutadiene can coating on a mild steel substrate, is shown in Fig. 1.8 [18]. The Fe2p3/2 spectrum of Fig. 1.8 is taken from the depth profile, acquired as described above, at the interface between oxide and polymer coating. The bulk of the oxide exhibited a Fe2p 3/2 spectrum which was entirely in the Fe(III) state (sketched in as a broken line on the leading edge of the spectrum of Fig. 1.8). At the interface Fe(II) character becomes visible in the spectrum, as a broadening at the lower binding energies (ca. 708 eV) and an Fe(II) shake-up satellite at ca. 716 eV. Thus one can assume that interfacial bonding between the polymer and the oxide involves at least a partial reduction of the Fe(III) surface oxide to Fe(II). This is not perhaps surprising as the polymer cures by an oxidative mechanism and at the interface, where there is a dearth of atmospheric oxygen, the iron oxide acts as the oxidizing agent for the cure process, being itself reduced to the lower valence state. It is conceivable that compound formation may take place with a product of the form of Fe(II) carboxylate. The result of such an interaction at the interphase is the formation of a discrete interphase zone, containing the reaction product, rather than a two-dimensional interfacial boundary between polymer and substrate. This was a rather novel con-

Fig. 1.7 Schematic diagram of the oxide stripping process. The metal substrate is dissolved in an appropriate solution (saturated iodine in methanol in the case of steel) leaving a duplex film of a thin oxide layer supported by the polymer. The integrity of the oxide surface is indicated by the SEM micrograph; sputter depth profiling can then be readily carried out through the oxide toward the polymer/metal interface.

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12


Fig. 1.8 Fe2p3/2 spectrum acquired at the interface between an iron substrate and a polybutadiene coating. The Fe(II) components result from the reduction of the iron oxide by the polymer.

cept at the time Ref. [18] was published, but now it has gained almost universal acceptance. While such chemical sectioning techniques have much to offer, taper sectioning and other mechanical removal methods have been developed to quite a high level of precision. This methodology was first applied in the semiconductor industry, where taper lapping of silicon wafers was used to expose a graded interface [19], and was then developed into a process known as ball-cratering [20], which used a large steel ball (ca. 30 mm in diameter) and fine diamond paste to erode a saucer-shaped depression, of controlled geometry, through the coating to the substrate, exposing the interfacial region. By measuring the diameter of the crater at any point (or the relative position of an analysis made by Auger electron spectroscopy) the position in depth relative to the original surface could be readily determined. This method was devised for hard coatings on metals and similar materials and was unsuitable for organic coatings such as paints. The solution turned out to be the modification of the stage of the ball-cratering machine to include cooling galleries to allow the passage of chilled nitrogen gas. By carefully controlling the temperature to a critical point below the Tg of the polymer, ball-cratering can be applied to polymer/metal interfaces [21]. The natural development of the taper sectioning and ball-cratering methodologies is to use a microtome to produce a well-defined section through the adhesive or coating in the same manner as is used to produce TEM sections. A method known as ultra-low-angle microtomy (ULAM) has recently been developed to enable very gentle tapers (0.03–28) to be prepared, and the interface to be analyzed by small-area XPS or ToF-SIMS. The methodology is fully described elsewhere [22] but in essence uses small blocks of stainless steel, that are out of


Fig. 1.9 Schematic of analysis procedure for depth profiling using ULAM sections. The table on the right shows typical depth resolutions achievable for small-area XPS at 15 and 100 lm, at a variety of angles.

parallel by small amounts to provide the taper, affixed to the bed of a histological microtome. When the taper has been established across the area of interest, the specimen is removed from the microtome and mounted for surface analysis. The basic geometry of the analysis process and the theoretical depth resolution that can be achieved are shown in Fig. 1.9. Using an X-ray spot of 15 lm in a small-area XPS spectrometer, a taper of 0.038 will yield a theoretical resolution of 13 nm. Data obtained in this manner from a polyvinylidene topcoat on a polyurethane primer are presented in Fig. 1.10; the depth scale has been recalibrated to relate the lateral distance along the specimen to the depth. It can be seen that the depth resolution achieved is of the order of that predicted by geometric calculations. ToF-SIMS analysis of the surface provides images of the interdiffusion of the two polymers and identifies the aggregation of a polyacrylic copolymer component of the PVdF coating at the interface [23]. The material and condition of the microtome knife are very important if one is to achieve a smear-free cut. Tungsten carbide knives have been found to produce excellent results on both thermosetting and thermoplastic polymers. Sectioning thickly coated metals is clearly not an option, but if thin (< 100 lm) metal foils are used as substrate material, it is quite possible to cut polymer/metal sections with a tungsten carbide knife. Model adhesive/substrate interfaces can usefully be prepared for microtoming and subsequent analysis from these foils of aluminum or iron, set into adhesives which are then cured. Fig. 1.11 shows a schematic of this concept which has been used for the study of epoxy/aluminum [24] and iron/polyamide interfaces [25].

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14


Fig. 1.10 Compositional depth profile through a PVdF (LHS)/ polyurethane (RHS) interface for a two-coat paint system acquired using ULAM preparation and 15 lm small-spot XPS [23].

Fig. 1.11 Schematic of model specimen prepared for the analysis of the adhesive/aluminum interface. The adhesive is simply cast around a sample of aluminum foil and then allowed to cure.

References

1.4 Conclusion

In the last few years there has been much progress in the ability to probe the interfacial chemistry of adhesion, as illustrated in the body of this paper. By the judicious use of model systems or careful chemical or mechanical sectioning there is now a reasonable expectation that such information should be accessible, given time, expertise, and the appropriate surface analytical instrumentation. One must not, however, lose sight of why such information is being sought: By developing a detailed knowledge of the manner in which the interfacial chemistry of a system influences its performance (durability), it should be possible to reverse-engineer the interface properties to provide the required systems performance. This is, indeed, the Holy Grail of those involved with design, specification, and selection of adhesives and organic coatings.

Acknowledgments

It is a pleasure to thank Professor Jim Castle and Drs Marie-Laure Abel, Steve Hinder, and Acharawan Rattana for their contributions to the work presented in this article.

References 1 J. F. Watts, J. Wolstenholme, An Introduc-

2

3

4 5

6

tion to Surface Analysis by XPS and AES, John Wiley, Chichester, 2003. D. Briggs, J. T. Grant, Surface Analysis by Auger and X-Ray Photoelectron Spectroscopy, IM Publications and SurfaceSpectra Ltd, Chichester and Manchester, 2003. J. Vickerman, D. Briggs, ToF-SIMS Surface Analysis by Mass Spectrometry, IM Publicationsand SurfaceSpectra Ltd, Chichester and Manchester, 2001. J. F. Watts, Surf. Interf. Anal., 12, 497– 503 (1988). J. F. Watts, in Handbook of Surface and Interface Analysis: Methods for Problem Solving (Eds: J. C. Riviere, S. Myhra), Marcel Dekker Inc., 1998, pp. 781–833. J. F. Watts, M.-L. Abel, in State-of-the-Art Application of Surface and Interface Analysis Methods to Environmental Materials Interactions (Eds.: D. R. Baer, C. R. Clayton, G. D. Davis, G. P. Halada), The Electro-

7 8 9 10 11

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chemical Society, Pennington, NJ, 2001, Vol. 2001–5, pp. 80–91. J. F. Watts, J. E. Castle, Int. J. Adhes. Adhes., 19, 435–443 (1999). J. E. Castle, R. Bailey, J. Mater. Sci., 12, 2049–2055 (1977). A. Rattana, M.-L. Abel, J. F. Watts, Int. J. Adhes. Adhes., in press (2005). S. R. Leadley, J. F. Watts, J. Adhes., 60, 175–196 (1997). J. F. Watts, S. R. Leadley, J. E. Castle, C. J. Blomfield, Langmuir, 16, 2292–2300 (2000). M.-L. Abel, I. W. Fletcher, R. P. Digby, J. F. Watts, Surf. Interf. Anal., 29, 115– 125 (2000). M. Sautrot, M.-L. Abel, J. F. Watts, J. Powell, J. Adhes., 81, 163–187 (2005). M. Guichenuy, J. F. Watts, M.-L. Abel, A. M. Brown, M. Audenaert, N. Amouroux, Surf. Interf. Anal., 36, 685–688 (2004).

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1 The Interfacial Chemistry of Adhesion: Novel Routes to the Holy Grail? 15 J. A. Bishopp, L. Davies, J. J. Haslam, Int. 16

17 18 19 20 21

J. Adhes. Adhes., 13, 11–119 (1993). A. Rattana, J. D. Hermes, M.-L. Abel, J. F. Watts, Int. J. Adhes. Adhes., 22, 205– 218 (2002). A. J. Kinloch, M. Little, J. F. Watts, Acta Materialia, 48, 4543–4553 (2000). J. F. Watts and J. E. Castle, J. Mater. Sci., 18, 2987–3003 (1983). M. L. Tarng, D. G. Fisher, J. Vac. Sci. Technol., 15, 50 (1978). J. M. Walls, D. D. Hall, D. E. Sykes, Surf. Interf. Anal., 1, 204 (1979). J. M. Cohen, J. E. Castle, Inst. Phys. Conf. Ser., Volume 93, Chapter 5, 275 (1988).

22 S. J. Hinder, C. Lowe, J. T. Maxted, J. F.

Watts, J. Mater. Sci., 40, 285–293 (2005). 23 S. J. Hinder, C. Lowe, J. T. Maxted, J. F.

Watts, Surf. Interf. Anal., 36, 1575–1581 (2004). 24 M.-L. Abel, J. F. Watts, A. Ottenwelter, J. Powell, Proc. 28th Annual Meeting of The Adhesion Society Inc., Feb. 13–16 2005, Mobile, AL, USA, pp. 155–157, 2005. 25 M. Guichenuy, S. J. Hinder, M.-L. Abel, J. F. Watts, in preparation.

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2 Modeling Fundamental Aspects of the Surface Chemistry of Oxides and their Interactions with Coupling Agents P. Schiffels, M. Amkreutz, A. T. Blumenau, T. Krüger, B. Schneider, T. Frauenheim, and O.-D. Hennemann

Abstract

Computer codes based on Density Functional Theory (DFT) and approximate Tight Binding (DFTB) have emerged as useful tools for the computational study of a wide range of materials in the last decade. The application of these tools to adhesion science is motivated by the fact that the surface chemistry of the substrates as well as mechanisms for substrate-adhesive interactions can be studied on an atomistic scale. In this chapter, we present recent developments which are useful for the study of molecular mechanisms relevant to adhesion. We first examine available theoretical methods for ideal silica surfaces and compare rigourous ab initio data to approximate DFTB results. It is shown that the latter method is sufficiently accurate for a detailed treatment of the systems. Furthermore, the Tight Binding scheme provides an efficient way to simulate the interactions of the organic components of the adhesive at the surface. In order to study these interactions, the DFTB approach is combined with a force field method in a novel QM/MM coupling scheme. Our work contributes to a molecular picture of the processes which build up the interphase between substrate and the polymer bulk of the adhesive.

2.1 Introduction: Atomistic Simulations in Adhesion

The detailed study of molecular mechanisms involved in adhesion requires an atomistic treatment of the substrate surfaces and their interaction with the organic components contained in the adhesive. Interesting aspects of the substrate–adhesive interaction include the preferential molecular orientation due to the interaction at the surface [1] or the influence of the initial stages of polymer grafting on the stability of polymer/metal interfaces [2]. The structure and composition of the interface can have a decisive effect on the properties of the reAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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2 Modeling Fundamental Aspects of the Surface Chemistry

sulting adhesive joint. Since the chemistry at the interface is rather complex, experimental studies on adhesion mechanisms often utilize surface analytical techniques in which the complexity is reduced by considering thin films, welldefined substrates, and model organic compounds [3]. The combined effort of experimental studies and atomistic simulations on model systems provides new insights into the surface chemistry relevant in adhesion [4]. In an applied field such as adhesion, these atomistic simulations have become possible only by the advent of efficient computer codes, and much work is thus based on density functional theory (DFT). Since these high-accuracy methods are computationally demanding if the size of the model system becomes large, appropriate approximate approaches are of the utmost importance for atomistic studies in the field of adhesion. An alternative to DFT is provided by tight binding schemes which have emerged as useful tools for the computational study of a wide range of materials in recent decades [5]. In this work, we focus on idealized model systems which are simulated using the density functional tight binding (DFTB) method, which is a second-order perturbation approach to the Kohn–Sham energy [6, 7]. DFTB results for substrate surfaces and organic species are compared with DFT results for validation purposes. We also present results obtained by a novel combined quantum mechanical/molecular mechanical (QM/MM) approach utilizing DFTB for the study of surface reactions. The latter coupling of accurate and approximate computational schemes becomes essential when considering reactions in more complex systems occurring in applied problems. QM/MM coupling leads to new perspectives for a fully atomistic treatment of surface–adhesive interactions in these complex systems because it can be used for systems comprising several hundreds of atoms. The present contribution is organized as follows. In Section 2.2, we describe DFTB results for ideal substrate surfaces and thus demonstrate the applicability of the approach for the new class of materials. We focus on silica, since silicon is one of the technologically relevant substrates encountered in microelectronic devices. In all practical applications, silicon is invariably covered by a large layer of amorphous silicon oxide. Atomistic investigations concerned with adhesive bonding on silicon therefore involve SiO2 surfaces and the specific chemistry of silica determines the surface reactions with the organic species present in adhesives. In order to validate the DFTB approach for silica surfaces, we predict reconstructions on ideal a-SiO2 (0001) surfaces by DFTB and compare the structures with those obtained by DFT methods. It is shown that geometrical parameters of the reconstructions as well as relative surface relaxation energies can be predicted accurately. Energetically favorable reconstructions on the (0001) a-SiO2 surface are presented which have not been described in the literature previously. These stable reconstructions are related to (SiO)4 and (SiO)6 rings with full coordination of both silicon and oxygen in the uppermost surface layer. The predicted stability of these rings makes them a good candidate for the atomistic structures of the cleaved surface. In Section 2.3 the DFTB approach is applied to the organic components of the adhesive and to model surface–adhesive interactions. In a first step, the pre-

2.2 Prediction of Surface Properties: Ideal Reconstructions on a-SiO2 (0001)

dicted geometries of important organic functional groups in epoxide and polyurethane adhesive systems are validated by DFT results. Afterwards, with the model for the silica substrate and the relevant organic species of the adhesive established, we move on to examine the direct interaction of coupling agents with the substrate and discuss the related chemical reactions. In general, significant computational effort can be saved if a high-accuracy method is restricted to local regions where high accuracy is essential, and a quantum mechanical treatment of the entire system is often unnecessary. QM/MM methods [8–10] are well suited if only a part of the system requires a high-level method; they have been applied to simulate reactions in solutions, enzymatic reactions and also adsorption processes [11]. In adhesion on silica, important reactions occur between the reactive functional groups of the adhesive and silanol groups at the surface. We examine reactions of silane coupling agents with silsesquioxane cage compounds which serve as model silica substrates. The restriction in size allows a comparison with DFT results feasible for systems of this size. It is shown that the DFTB treatment of the direct chemical interaction between the coupling agent and the model silica substrate compares well with DFT results and predicts the energetical order of the reactions energies correctly.


The surface chemistry of amorphous silica has been studied experimentally in some detail in recent decades [12, 13]. Previous theoretical studies on these surfaces employed demanding either ab-initio Car–Parrinello dynamics [14] or molecular dynamics [15, 16] techniques which are based on empirical potentials. For the investigation of the growth mechanism of amorphous SiO2 layers on Si(001), a modified Stillinger-Weber potential was developed [17, 18]. Some studies were also aimed at identification of reactive surface sites via cluster calculations [19], since these sites are relevant for several technological processes. Furthermore, hydroxylation reactions occur preferentially depending on the specific site. In order to test the applicability of the DFTB approach for silica surfaces, we investigate one of the morphologically stable surfaces of ideal a-quartz – the (0001) surface – in detail. In principle, the structure of this surface depends on the atmospheric conditions, and especially on the presence of adsorbed water due to atmospheric humidity. For simplicity, we focus on surfaces obtained in perfect vacuum. Ideal a-quartz surfaces have been the subject of a number of recent studies using either first principles [20, 21] or empirical potential approaches [22, 23]. The reconstructions of the a-SiO2 (0001) surface are obtained using three-dimensional periodically repeated supercells including a sufficiently large vacuum region to inhibit interaction between opposite surfaces of a slab and with the periodic images of the slab. One surface of the slab was terminated by a complete H monolayer in order to saturate dangling bonds of the silicon atoms and

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to minimize the charge transfer between the surfaces, thus leading to a better representation of the bulk material. The induced Mulliken charges near the hydrogen-terminated surface indicate that the effect is negligible after a few bonds. Atomic positions close to the hydrogen-passivated end of the model are constrained to the respective bulk positions in order to preserve the ideal bulk configuration. The DFTB results for the surface reconstructions are based on the C point approximation in the Brillouin-zone sums. Details are given in Ref. [5]. All reconstructions are modeled in a hexagonal (2 ´ 2 ´ 3) cell with three times the period length along [0001] – resulting in eight silicon layers. It was checked that the structures obtained converged with respect to slab thickness. Generally, the bulk structure is observed to recover within 5 Å from the surface, showing that the slab considered is sufficiently thick for a representation of the semi-infinite solid. Analogous DFT calculations were performed within the local density approximation (LDA) employing the plane-wave code CASTEP [24] with the parameterization by Perdew and Zunger (CA-PZ) [25, 26]. These plane-wave calculations use the standard Vanderbildt ultra-soft pseudopotentials [27] from the CASTEP database. We found that a cut-off energy of 380 eV for the plane-wave basis is sufficient to obtain converged structures and relative energies even though the use of ultra-soft pseudopotentials for silicates has been a matter of debate recently [28]. For calculations within the generalized gradient approximation (GGA), we employed the DMol3 program [29, 30] and the BLYP parameterization [31, 32] with a numerical double-f basis plus polarization functions (DNP). It had been observed earlier that the energetics and phase stabilities for the SiO2 polymorphs depend strongly on the type of approximation used in density functional theory [33, 34]. For the investigation of the surface reconstructions, we decided on BLYP and CA-PZ since in our calculations these functionals reproduce the correct order of relative stabilities for silica polymorphs. All DFT calculations on the surface reconstructions were done with slightly thinner slabs of five layers of silicon, while Brillouin-zone sums were performed in a uniform Monkhorst-Pack grid including five special k-points in both cases. Initial structures for the optimizations were generated manually, followed by relaxation with respect to internal coordinates while the surface net parameters are fixed to those obtained for the optimized a-quartz structures (a = b = 9.92 Å (DFTB); 9.84 Å (CA-PZ); 10.04 Å (BLYP)). The resulting stable surface reconstructions for the SiO-terminated surface are presented in Fig. 2.1. In a-quartz, all layers of SiO termination along [0001] are equivalent, while oxygen-terminated surfaces have been found to be unstable. The first SiO-terminated surface 1 was obtained by simply cleaving the bulk material and is referred to as the “cleaved surface” in Ref. [21]. The formation of Si–O–Si bridges with all atoms in bulk coordination is energetically unfavorable as this would lead to high bond length and angle distortions. In the case of the (2 ´ 1) surface reconstruction 2 – the “VAP surface” [21] – all silicon atoms are fully coordinated and one valence alternation pair is formed per (2 ´ 1) unit in which there is a onefold and a threefold coordinated oxygen center. The (2 ´ 1) reconstruction 3 has not been described previously in


Fig. 2.1 Reconstructions of the clean a-quartz (0001) surfaces obtained by the BLYP method. The structures are presented as top views along the [000-1] direction. Oxygen atoms are represented by dark, silicon atoms by light colour coding. The topmost layer is shown in a ball-and-stick representation, for clarity.

the literature. It consists of rows of parallel Si–O–Si bridges, leaving all atoms fully coordinated in a corrugated surface. In the final relaxed structure the two topmost Si layers are combined into one layer, giving a surface of higher density than the preceding structures. In reconstruction 4, two (SiO)3 rings per (2 ´ 1) surface unit are aligned in a chain connected by Si–O–Si bridges. The structure is labeled “semi-dense” according to Ref. [21], since the two topmost silicon layers are only partially merged into one layer. Also, in this case once again all atoms have full bulk-like coordination. The last two (1 ´ 1) reconstructions studied – 5 and 6 – are related in the sense that both contain (SiO)6 rings in the topmost layer. The latter reconstructions lead to Si–O–Si bridges which – in contrast to 3 – are not parallel, but are arranged in ring-like structures. It is interesting to note that reconstructions containing (SiO)2 rings were found to be energetically unstable. Structures showing a considerable number of under-

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coordinated atoms can be expected to be highly reactive, whereas we found two reconstructions on the (0001) surface that have not been described previously – 3 and 5 – with full coordination of both silicon and oxygen in the uppermost surface layer. The bulk-like coordination of both oxygen and silicon centres results in particularly low surface energies, as can be inferred from Table 2.1. The predicted stability of the (SiO)6 rings makes them good candidates for the atomistic structures of the cleaved surface in vacuum. Regarding the applicability of the DFTB approach for the prediction of surface properties, we note that the relative ordering of stabilities is predicted correctly by DFTB. The results compare well with our own DFT results and as well as with the values obtained in the literature [21, 22]. Deviations in the absolute values for the surface relaxation energies between DFTB and BLYP are of the same order of magnitude as the deviations between the two DFT calculations (CA-PZ and BLYP). Concerning structural data, we compare the (SiO)6 rings of the lowest-energy reconstruction 6, which are in excellent agreement with data available in the literature.

Table 2.1 Surface relaxation energies for a-SiO2 (0001): computed surface relaxation energies of the surface models investigated for the clean a-quartz (0001) surface. a)

Reconstruction

Description

Property b)

1 (1 ´ 1)

“cleaved”

2 (2 ´ 1)

“VAP”

CA-PZ

BLYP

DFTB

Erelax Erel

–61 130

–52 97

–76 87

120

124

Erelax Erel

–97 94

–63 86

–92 71

90

49

3 (2 ´ 1)

(SiO)4 rings

Erelax Erel

–122 69

–71 78

–108 55

4 (2 ´ 1)

“Semidense” (SiO)3 rings

Erelax Erel

–131 59

–86 63

–122 41

5 (1 ´ 1)

(SiO)6 rings

Erelax Erel

–164 27

–121 28

–138 25

6 (1 ´ 1)

“Dense” (SiO)6 rings

Erelax Erel

–191 0

–149 0

–163 0

122 126 134

120 125 131

122 119 128

Si–O–Si angles [deg] a) b)

Selected structural parameters have been compared with available literature data. Erelax refers to the relaxation energy, while Erel gives the relative energies with respect to the lowest-energy surface 6. All energies are given in meV Å–2.

Ref. [21] Ref. [22]

40

0

0

123 123 129

2.3 Organic Components of the Adhesive and Substrate-Adhesive Interaction


In this section, we apply the density functional tight binding approach to the organic components of epoxide and polyurethane adhesive systems and study the interaction of silane coupling agents with model silica substrates. The quality of the DFTB results for the organic compounds is first validated by comparison of predicted geometrical parameters with DFT results. To this end, we define a test set of organic compounds containing functional groups relevant for epoxides and polyurethanes. The test set includes ethylene oxide, formic acid, ethylene glycol, and 1-hydroxy-2-ethylamine as model compounds for epoxides as well as isocyanate, carbamic acid, and methyl carbamate for the polyurethane adhesive systems. In addition, we include diethylenetetramine, maleic acid anhydride, the diglycidyl ether of bisphenol A, tripropyleneglycol and diphenylenemethane 4,4'-diisocyanate in the present study. Important coupling agents in commercial application are (3-glycidyloxypropyl)trimethoxysilane and triethoxyvinylsilane. As for the coupling agents, we consider the products of the hydrolysis reaction rather than the corresponding alkoxysilanes because the former are assumed to be the reactive species at the substrate. These are abbreviated to GOTHS for the (3-glycidyloxypropyl)silane and as THVS for the vinylsilane. In Table 2.2, we present results for selected bond lengths and angles of the two coupling agents predicted by DFTB and compare these with corresponding DFT structures. The all-electron DFT results used as a reference are carried out with the DMol3 program [29, 30] using a numerical double-f basis plus polarization functions (DNP). We employ both LDA (VWN) as well as GGA calculations (BLYP, PW91) and have obtained structures which converge to a maximum dis-

Table 2.2 Structural parameters for GOTHS and THVS: selected structural parameters calculated using the different DFT functionals and the DFTB method for the two silane coupling agents studied in this work.

System

Geometrical parameter

BLYP

PW91

VWN

DFTB

GOTHS

r(SiO) r(SiC) r(OH) r(CC)epoxide (HCH)epoxide (COC)epoxide

1.66 Å 1.88 Å 0.97 Å 1.47 Å 115.48 61.08

1.65 Å 1.86 Å 0.96 Å 1.47 Å 115.58 61.38

1.63 Å 1.84 Å 0.97 Å 1.46 Å 115.78 61.98

1.63 Å 1.87 Å 0.97 Å 1.47 Å 115.28 61.38

THVS

r(C = C) r(SiO) r(SiC) r(OH) (HCH)

1.34 Å 1.65 Å 1.86 Å 0.97 Å 115.88

1.34 Å 1.65 Å 1.85 Å 0.96 Å 116.68

1.33 Å 1.63 Å 1.83 Å 0.97 Å 117.28

1.33 Å 1.62 Å 1.85 Å 0.96 Å 115.58

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24


placement of 0.0025 Å. Details concerning the DFTB calculations and formalism are given in Ref. [5]. The geometrical data for the two silanes obtained by DFTB are in excellent agreement with the computationally more expensive DFT results, as can be inferred from Table 2.2. This statement can be generalized to other organic species of the test set (Fig. 2.2). The mean absolute deviation between BLYP and DFTB structures for the whole test set is only 1.1% for the bond lengths and 1.9% for the corresponding angles. With the DFTB model for the silica substrate and the organic components established, we now focus on the reaction of coupling agents with model silica substrates. As mentioned in the introduction (Section 2.1), for these surface–adhesive interactions we follow a novel QM/MM approach in which DFTB serves as a QM method which is combined with a molecular mechanics (MM) approach. The QM/MM coupling is affected by means of the ONIOM2 scheme as introduced by Morokuma and co-workers [10, 35–39] and implemented using the Gaussian 03 program suite [40]. The general set-up of the method is depicted in Fig. 2.3. In the ONIOM2 embedding scheme, the simulated system is generally divided into an inner model system and a surrounding region. The inner model system – which is usually the reactive part of the system – is defined by cutting bonds between the inner system and the surroundings, which are saturated with hydrogen link atoms. In our implementation, we treat the inner model system with DFTB and the remaining atoms are described by the universal force field (UFF) [41] in the parameterization provided by Gaussian 03. This force field has already been applied in QM/MM studies of adsorption reactions on c-alumina [11]. As shown in Fig. 2.3, the desired energy of the whole system using the high-level method is approximated by two contributions: the first is the energy obtained for the inner model system using the high-level approach; the second is approximated by the energy difference DE between the inner model system and the whole system obtained using the low-level method. A projection of forces has to be done to take into account the interaction of the different layers. Details can be found in Refs. [10, 35–39]. The coupling ensures the necessary accuracy of the DFTB scheme for the direct interaction region, while compromising on accuracy for the rest of the model system for efficiency reasons. The QM(DFTB)/MM(UFF) approach is generally capable of treating reactions between adhesives and model surfaces consisting of several hundreds of atoms in either periodic supercells or as cluster models. For validation purposes, we chose silsesquioxanes as well-defined models for the hydroxylated silica surfaces. Silsesquioxanes [42] are silicon- and oxygen-containing cage compounds with the general stoichiometry [RSiO3/2]n which are built of edge-sharing [RSiO3] tetrahedra. We introduce the abbreviation TnRn where T represents the [SiO3/2] unit and R indicates some substituent. In this work, we focus on cages with one silanol group and hydrogen termination (TnHn–1(OH) with n = 4, 6, 8). The smaller systems consist of 20 to 60 atoms and thus can also be treated by DFT. Equations (1–3) represent possible reactions of the two coupling agents – GOTHS and THVS – with the silanol group of the silsesquioxanes. For


Fig. 2.2 Correlation diagrams illustrating the accuracy of the DFTB approach for typical organic functional groups relevant for epoxide and polyurethane adhesive systems. Structural parameters (top: distances;

bottom: angles) obtained by DFT (VWN, PW91) and DFTB are compared with BLYP results, which are taken as the reference. The test set of organic molecules is defined in the text.

GOTHS we consider two competing reactions for the substituted silsesquioxanes TnHn–1(OH) with n = 4, 6, 8 – the condensation [Eq. (2)] and the cleavage of the oxirane ring [Eq. (3)]. For THVS, we consider the condensation [Eq. (1)] only. BLYP structures of the reaction products for the latter reaction are presented in Fig. 2.4. In the QM/MM calculations, the fragment [O3Si-O-SiO2C] is

25

26


Fig. 2.3 Concept of the two-layer ONIOM approach (ONIOM2) used for the QM/MM coupling in this work, connecting a high-level method – in our case DFTB – with a lowlevel method, in our case a force field approach.

Fig. 2.4 Structures predicted by BLYP for the products of the THVS condensation to the silsesquioxanes TnHn–1(OH) with n = 4, 6, 8.


defined as the QM(DFTB) zone for the reaction products of Eqs. (1) and (2). In the case of Eq. (3), the fragment [O3Si–O–CHCH2OH] is considered as a QM(DFTB) zone.

In order to check the accuracy of the QM(DFTB)/MM(UFF) approach, we first compare the geometries of the reaction products with BLYP results obtained with Gaussian 03 (cc-pVDZ basis set). The results are discussed considering the oxirane cleavage reaction [Eq. (3)] as a typical example. Generally, it is observed that structural deviations between QM/MM and BLYP for the reactions involving the T4H3(OH) cluster are negligible. This may be due to the small cluster size leading to a more rigid structure with less sensitivity to distortions. Considering the reaction products with the larger clusters, T6H5(OH) and T8H7(OH), some minor deviations in the dihedral angles for the coupling agent are observed, while the structures of the silsesquioxane cages differ considerably. In these structures, SiO4 tetrahedra assigned to the MM zone are distorted compared with the BLYP structures. In detail, this behavior can be seen in Fig. 2.5, where the QM/MM structure for the reaction product of GOTHS and T8H7(OH) from Eq. (3) are compared with BLYP results. The bond lengths agree very well 7 within a maximum deviation of 5% 7 compared with BLYP but the bond angles show noticeable discrepancies with regard to the BLYP data in the range 145–1508. These deviations are related to an underestimation of the Si–O–Si angles by QM/MM introduced by the universal force field. This is confirmed by pure DFTB calculations, where noticeable improvements are observed (Fig. 2.5). The distortion of the Si–O–Si angles in the T8H7(OH) cluster disappears and the mean absolute deviation of the angles from the BLYP results is reduced to less than 5%. In a second step, we apply the QM(DFTB)/MM(UFF) approach to compare the total electronic energy differences for the nine reactions of Eqs. (1–3). The

27

28


Fig. 2.5 Comparison of selected structural parameters predicted by the QM(DFTB)/ MM(UFF) scheme (top) and the corresponding DFTB results (bottom) for the product of Eq. (3) (oxirane cleavage reaction of

GOTHS at the T8H7(OH) cluster). Bond lengths [Å] and bond angles [deg] are compared with the results of BLYP calculations which serve as a reference.

2.4 Conclusion and Outlook Table 2.3 Reaction energies for the reactions considered in Eqs. (1–3): total electonic energy differences [kcal mol–1] including zero-point vibrational corrections for the reactions of THVS and GOTHS with the silsesquioxanes.

Reaction

ONIOM2: QM(DFTB)/MM(UFF)

BLYP

Eq. (1) n=4 n=6 n=8

T4H3(OH) + THVS T6H5(OH) + THVS T8H7(OH) + THVS

–2.8 –3.0 –4.0

–1.7 +0.1 –1.2

Eq. (2) n=4 n=6 n=8

T4H3(OH) + GOTHS, condensation T6H5(OH) + GOTHS, condensation T8H7(OH) + GOTHS, condensation

–8.9 –7.8 –4.9

–2.7 –1.5 –0.5

Eq. (3) n=4 n=6 n=8

T4H3(OH) + GOTHS, cleavage T6H5(OH) + GOTHS, cleavage T8H7(OH) + GOTHS, cleavage

–36.7 –38.1 –39.1

–22.6 –22.6 –22.7

educts and products are simulated using the QM/MM method, ensuring that the same atoms are chosen for the inner QM zones of the educts and the QM zone of the product of a reaction, respectively. Table 2.3 gives the reaction energies obtained as the difference of total educt and product electronic energies in ONIOM2 – corrected by their respective zero-point vibration contributions – as well as the corresponding BLYP results. The results of the ONIOM2 approach are in moderate agreement with the DFT data concerning the absolute values for the reaction energies. Nevertheless, the ranking of the reactions in terms of energy is predicted correctly by QM/MM compared with DFT, which makes the QM/MM coupling a predictive tool. In all cases, the reaction energies for the condensation reactions are of the same order of magnitude, while the oxirane cleavage is definitively energetically favored. This behavior is described very well by the application of the QM/MM approach, in agreement with the DFT results.

2.4 Conclusion and Outlook

In this work, we applied the DFTB method to some fundamental aspects of adhesion on oxide surfaces and validated the approach by comparison with rigorous DFT calculations. In Section 2.2, DFTB results for surface reconstructions on ideal a-SiO2 were investigated to serve as a benchmark for surface structures. In Section 2.3, we presented calculations concerning the structures of organic components of the adhesive and introduced a new computational scheme based on QM/MM coupling for surface-adhesive interactions.

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30


The investigations on surface reconstructions on ideal a-SiO2 produced two structures, in addition to those described in the literature, which are related to stable (SiO)n rings in the topmost surface layer. The DFTB results for these reconstructions show convincing agreement on the relative order of surface relaxation energies as compared with DFT. The agreement of structural parameters with available literature data is excellent. For the organic components of the epoxide and polyurethane adhesive systems, we investigated the structural parameters predicted for a test set of organic species with the relevant functional groups. The mean absolute deviation between the BLYP and DFTB structures is only 1.1% for the bond lengths and 1.9% for the corresponding angles. Hence the DFTB method appears to be an adequate approximate tool for the simulation of organic components and surface structures which are relevant for the adhesive bonding on silica. For the direct interaction of coupling agents with silica substrates, we introduced a novel QM/MM scheme. This scheme is based on the ONIOM2 approach which combines the DFTB method with a force field (UFF). The scheme was applied to reactions between coupling agents and silsesquioxanes which serve as well-defined models for the hydroxylated silica surface. It could be shown that the QM/MM geometries of the reaction products can be predicted reasonably well by the QM(DFTB)/MM(UFF) approach and that the quality of the prediction depends on the size of the inner QM zone simulated by DFTB. In the case of reaction energies, a moderate agreement compared with DFT is found, but the order of reaction energies and thus the prediction of the preference of a reaction is described well. It was found that the oxirane cleavage reaction of GOTHS is energetically favored compared with the condensation. The prediction of reaction energies by QM(DFTB)/MM(UFF) is at most as reliable as the accuracy achievable by the high-level method in the coupling scheme – in this case DFTB. In order to improve the results for the reaction energies, a combination of DFTB with a DFT method is desirable. In such a QM(DFT)/QM(DFTB) scheme, the reactive functional groups can be treated by DFT, while the DFTB method is used for the description of the rest of the system. The QM(DFT)/QM(DFTB) coupling can be used for larger systems than pure DFT approaches, due to the approximate character of the DFTB method. Finally, a QM/QM/MM scheme could be applied to simulate even larger systems by using a force field method which is useful according to the results shown in this work.

References 1 S. Dieckhoff, R. Hoeper, V. Schlett, T.

Gesang, W. Possart, O.-D. Hennemann, J. Guenster, V. Kempter, Fresenius’ J. Anal. Chem. (Germany) 1997, 358, 258– 262.

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Journal of Adhesion and Adhesives 1999, 19, 425–434. B. Schneider, O.-D. Hennemann, W. Possart, Journal of Adhesion 2002, 78, 779. T. Frauenheim, G. Seifert, M. Elstner, T. Niehaus, C. Kohler, M. Amkreutz, M. Sternberg, Z. Hajnal, A. Di Carlo, S. Suhai, Journal of Physics: Condensed Matter 2002, 14, 3015–3047. M. Elstner, D. Porezag, G. Jungnickel, J. Elsner, M. Haugk, T. Frauenheim, S. Suhai, G. Seifert, Physical Review B (Condensed Matter) 1998, 58, 7260–7268. M. Elstner, T. Frauenheim, E. Kaxiras, G. Seifert, S. Suhai, Physica Status Solidi B 2000, 217, 357–376. D. Bakowies, W. Thiel, Journal of Physical Chemistry 1996, 100, 10 580–10 594. M. J. Field, P. A. Bash, M. Karplus, Journal of Computational Chemistry 1990, 11, 700–733.

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2 Modeling Fundamental Aspects of the Surface Chemistry K. N. Kudin, J. C. Burant, J. M. Millam, S. S. Iyengar, J. Tomasi, V. Barone, B. Mennucci, M. Cossi, G. Scalmani, N. Rega, G. A. Petersson, H. Nakatsuji, M. Hada, M. Ehara, K. Toyota, R. Fukuda, J. Hasegawa, M. Ishida, T. Nakajima, Y. Honda, O. Kitao, H. Nakai, M. Klene, X. Li, J. E. Knox, H. P. Hratchian, J. B. Cross, C. Adamo, J. Jaramillo, R. Gomperts, R. E. Stratmann, O. Yazyev, A. J. Austin, R. Cammi, C. Pomelli, J. W. Ochterski, P. Y. Ayala, K. Morokuma, G. A. Voth, P. Salvador, J. J. Dannenberg, V. G. Zakrzewski, S. Dapprich, A. D. Daniels, M. C. Strain, O. Farkas, D. K. Malick, A. D. Rabuck, K. Raghavachari, J. B.

Foresman, J. V. Ortiz, Q. Cui, A. G. Baboul, S. Clifford, J. Cioslowski, B. B. Stefanov, G. Liu, A. Liashenko, P. Piskorz, I. Komaromi, R. L. Martin, D. J. Fox, T. Keith, M. A. Al-Laham, C. Y. Peng, A. Nanayakkara, M. Challacombe, P. M. W. Gill, B. Johnson, W. Chen, M. W. Wong, C. Gonzalez, J. A. Pople, Gaussian 03, Revision B.04, Gaussian, Inc., Pittsburgh, PA, 2003. 41 A. K. Rappe, C. J. Casewit, K. S. Colwell, W. A. Goddard, III, W. M. Skiff, Journal of the American Chemical Society 1992, 114, 10 024–10 035. 42 M. G. Voronkov, V. J. Lavrentyev, Topics Curr. Chem. 1982, 102, 199.

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3 Adhesion at the Nanoscale: an Approach by AFM M. Brogly, O. Noel, G. Castelein, and J. Schultz

Abstract

The atomic force microscope (AFM) is a promising device for the investigation of materials surface properties at the nanoscale. Precise analysis of adhesive and mechanical properties, and particularly of model polymer surfaces, can be achieved with a nanometer probe. This study distinguishes the different contributions (chemical and mechanical) included in an AFM force–distance curve in order to establish relationships between interfacial tip–polymer interactions and the surface viscoelastic properties of the polymer. Measurements have been made, in the AFM contact mode, of both chemical and mechanical local attractive or adhesive forces of model substrates. Assuming that the main technical uncertainties have been listed and minimized, surface force measurements were first performed on chemically modified silicon substrates (grafted with hydroxyl, amine, methyl, and ester functional groups). The surface chemistry contribution (in particular, its hydrophilic features) is dominant in the measurement of the adhesion force. A linear relationship has been obtained between the van der Waals component of the thermodynamic work of adhesion and the surface energy of the silicon grafted substrates. To investigate the effects of mechanical or viscoelastic contributions, we made force measurements on model polymer networks whose surfaces were chemically controlled with the same functional groups as before (silicon substrates). Young’s modulus has been determined on the basis of nanoindentation experiments. Only the beginning of the loading regime was considered, as Hertz theory could then be applied. The viscoelastic contribution was dominant in the adhesion force measurement. We propose new relationships between the local adhesion force, the dissipation energy in the tip–polymer contact, and the surface properties of the material (thermodynamic work of adhesion), and show that the dissipation function is related to Mc, the mass between crosslinks of the network.


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3 Adhesion at the Nanoscale: an Approach by AFM

3.1 Introduction

The atomic force microscope (AFM) is a promising device for the investigation of the surface properties of materials at the nanoscale. Precise analysis of adhesive and mechanical properties, and in particular of model polymer surfaces, can be achieved with a nanometer probe. The purpose of this study is to distinguishes the different contributions (chemical and mechanical) included in an AFM force–distance curve in order to establish relationships between the interfacial tip–polymer interactions and the surface viscoelastic properties of the polymer. Indeed, we are interested in the measurements of local attractive or adhesive forces, in AFM contact mode, of model substrates, from both the chemical and mechanical viewpoints. These measurements are performed with a silicon nitride (Si3N4) tip, at ambient temperature, in the air. With the assumption that the main technical uncertainties have been listed and minimized, surface force measurements are, in a first step, performed on chemically modified silicon substrates (grafted with hydroxyl, amine, methyl, and ester functional groups). In order to investigate the effects of mechanical or viscoelastic contributions, we achieved force measurements on model polymer networks, whose surfaces are chemically controlled with the same functional groups as before (on silicon substrates). Young’s modulus has been determined on the basis of nano-indentation experiments. Only the beginning of the loading regime was considered, as Hertz theory could then be applied. The results show that the viscoelastic contribution is dominant in the adhesion force measurement. Finally, we propose new relationships which express the local adhesion force to the dissipation energy in the tip–polymer contact and the surface properties of the material (thermodynamic work of adhesion). Moreover, we show that the dissipation function is related to Mc, the mass between crosslinks of the network.

3.2 Materials and Methods

Si(100) silicon wafers (Mat Technology, France) polished on one side were used as the substrate for self-assembled monolayer (SAM) film grafting. In this paper, “as received silicon” (Sias received) refers to a silicon wafer previously cleaned with ethanol in an ultrasonic bath. That means that a contaminated layer still remains on the surface. Four organosilane grafts (supplied by ABCR, Karlsruhe, Germany) were used for the elaboration of homogeneous model surfaces on the substrate. Two hydrophobic model surfaces were prepared by using hexadecyltrichlorosilane (C16H42O3Si or SiCH3) and 1H,1H,2H,2H-perfluorodecylmethyldichlorosilane (C11H7Cl2F17Si or SiCF3) and two hydrophilic model surfaces by using (6-aminohexyl)aminopropyltrimethoxysilane (C12H30N2O3Si or SiNH2) and (2-carbomethoxy)ethyltrichlorosilane (C4H7Cl3O2Si or Siester). Polydimethylsiloxane (PDMS) was supplied by ABCR (Karlsruhe, Germany). All other


chemicals used in chemical handling (cleaning, synthesis) were of reagent grade or better (supplied by Aldrich). 3.2.1 Preparation of Oxidized Silica Surface

Before coating, the substrates must be chemically modified in order to obtain a hydrophilic surface (SiO2). The silicon surface is first cleaned with ethanol and dried with nitrogen before oxidation. Oxidized surfaces are obtained after cleaning the substrate in a warm Piranha (60 8C) solution (30% H2O2/H2SO4 mixture, 3:7 v/v) for about 30 min in order to keep a smooth surface, and then thoroughly rinsing with deionized and twice-distilled water. Just before being grafted with organosilane the wafers are dried with nitrogen. This treatment produces a high hydroxyl group density on the surface (SiOH groups), to which functional silanes will be adsorbed upon hydrolysis [1]. Wafers of silicon covered with hydroxyl end-groups (SiOH) were synthesized by this method and immediately probed in order to avoid contamination of the surface by the environment due to the high reactivity of SiOH groups. 3.2.2 Grafting of Functionalized SAMs onto Silicon Wafer

Three different techniques are frequently used to obtain SAMs: Langmuir–Blodgett techniques, involving an air–water interface to transfer the assembled film to a solid substrate; solution adsorption of film molecules onto the substrate; and a vapor-phase molecular self-assembling technique [2], which uses vapor deposition of the film onto the substrate. Our functionalized SAMs were prepared by the last of these techniques, which had been slightly improved in the laboratory [1]. The lack of solvent prevents possible incorporation of small solvent molecules contamination and defects into the SAMs. Moreover, a previous study [3] showed that the molecular films prepared by this method are more homogeneous. The silicon wafers were placed above a previously de-aired solution consisting of a mixture of 100 lL organosilane and 3 mL paraffin. The vapor-phase deposition of the molecular film on the substrate was performed in a vacuum chamber (50 min at 5 ´ 10–3 Torr) at room temperature. 3.2.3 Crosslinking and Functionalization of PDMS Networks

PDMS samples were crosslinked under nitrogen in a glovebox using tetrakis(dimethylsiloxy)silane as a crosslinker, and a platinum-based catalyst. All the chemicals were supplied by ABCR (Karlsruhe, Germany). The classification of PDMS substrates refers to the length of the chains before crosslinking (determined by SEC) (Table 3.1). Then, PDMS 1.5k is the hardest substrate, whereas 53k refers to the softest one. Flory’s law of rubber elasticity, which represents

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3 Adhesion at the Nanoscale: an Approach by AFM Table 3.1 Mechanical properties of the crosslinked PDMS.

PDMS

Initial chain mass [g mol–1]

Elasticity domain [%]

Deformation at break [%]

Young’s modulus [MPa]

1.5k 7k 11k 53k

1 500 7 000 11 000 53 000

40 46 47 52

196 210 250 250

2.24 0.64 0.30 0.13

Fig. 3.1 Variation of 1/E as a function of Mc [f (Mc) curve], where E is the Young’s tensile modulus determined with a 1 mm min–1 strain rate, and Mc [g mol–1] is the average mass between two crosslinks.

the reciprocal tensile modulus versus Mc, the mass between crosslinks, is satisfied for the networks synthesized, as reported in Fig. 3.1. This proves that a good control of macroscopic mechanical properties is achieved. Moreover, the Young’s moduli are independent of the strain rate in the range of strain rates corresponding to those used for AFM experiments and at ambient temperature for all the substrates. PDMS network substrates are then treated with water plasma and functionalized by the vapor deposition technique. 3.2.4 Characterization of the SAMs

AFM topographic images, X-ray photoelectron spectroscopic (XPS) analysis, and ellipsometric and contact angle with water droplet (Table 3.2) measurements

Table 3.2 Contact angle of water droplets on SAMs of various terminal functionalities.

SAM functionality

Water contact angle he [deg]

CF3 CH3 Sias received COOR NH2 OH

106 103 78 71 57 6

3.3 Results and Discussion

were performed on the SAMs and showed that we had obtained homogeneous, well-packed, organized, and stable grafting on both the silicon wafers and PDMS.

3.3 Results and Discussion 3.3.1 Force–Distance Curve Measurements and AFM Calibration 3.3.1.1 Force–Distance Curve Features Force measurements with an AFM in the contact mode consist in detection of the deflection of a spring (or cantilever), bearing a silicon nitride tip, when interacting with the sample surface. The deflection of the cantilever is detected by an optical device (four-quadrant photodiode) while the tip is vertically moved forward and backward thanks to a piezoelectric ceramic (or actuator). Thus, provided that the spring constant of the cantilever is known, one can obtain a deflection–distance (DD) curve and hence a force–distance (FD) curve, by applying Hooke’s law. The DD curves were obtained in air with commercially available apparatus (Nanoscope IIIa D3000; Digital Instruments). Fig. 3.2 is a schematic representation of a DD curve obtained when probing a hard surface.

3.3.1.2 The DD Curve (Contact Mode) In Fig. 3.2, zone A, the cantilever is far from the surface and stays in a state of equilibrium (no interaction with the surface). The cantilever deflection is zero. During the approach toward (or withdrawal from) the surface, the tip interacts

Fig. 3.2 Schematic representation of a DD curve. The slope of the “contact zone” is equal to unity when considering the contact between the tip and a rigid surface, whereas it is lower than 1 for the contact between the tip and a soft material (polymer).

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with the sample and a jump-in (or jump-off) contact occurs (zones B (for loading) and E (for unloading)). These instabilities take place because the cantilever becomes mechanically unstable. Usually, for undeformable surfaces, because of mechanical instabilities, a jump-in contact is too small to determine attractive van der Waals forces. When in contact, the cantilever deflection is equal to the piezoelectric ceramic displacement, provided no indentation of the substrate occurs (zones C (for loading) and D (for unloading)). An undeformable reference sample (a cleaned silicon wafer) is used to scale the DD curve in deflection by fixing the value of the slope of the contact line at unity.

3.3.1.3 AFM Calibration One of the fundamental steps to obtaining reproducible, quantitative, and reliable data is the calibration procedure, which should be rigorous and systematic for all measurements.

3.3.1.3.1 Determination of the Spring Constant of the Cantilever For this determination we focused on a nondestructive method, based on the use of reference rectangular cantilevers [4]. The cantilever used in this study was a triangular cantilever (supplied by Nanosensor, Germany) with an effective 0.30 ± 0.03 N m–1 spring constant (the one specified by the supplier is 0.58 N m–1!).

3.3.1.3.2 Nonlinearity of the Quadrant of Photodiodes The nonlinearity of the optical detector is the consequence of an inhomogeneous spreading of the laser spot on the detector. This nonlinearity has been studied by reporting the slope of the contact line (zones C or D) of a DD curve (obtained on a hard surface and assuming that there is no nonlinearity at the middle of the photodetector) versus the tension (volts) measured by the detector. The domain of linearity of the detector appeared to lie between +2 V and –2 V. If nonlinearity is not taken into account the error on the quantitative results can be significant, because the slope of the contact line determines the Y-scale.

3.3.1.3.3 Scan Rate of the Cantilever The actuator shows a hysteresis in its vertical displacement. This hysteresis can be studied by reporting the slope of the contact zones (zones C and D) versus the amplitude of the contact zone and the scan rate. During the experiments, the actuator is considered to be thermally stable. We observed that a discrepancy appears for very low scan rates. For higher scan rates, the viscosity of the environment could be significant. A rate of about 6 lm s–1 is a good compromise.


3.3.1.3.4 Systematic Check In addition, checking the adhesion force on a reference silicon wafer regularly and randomly monitors contamination of the tip during the measurements. When the tip is contaminated, a new tip is used and characterized. In that way, we can select tips with about the same radius and the same spring constant in order to compare the experimental values. Tip contamination occurs rarely in comparison with the large number of DD curves that are realized. Finally, the results are reported as an average of about 100 DD curves for each substrate. 3.3.2 Force–Distance Curves on Rigid Systems of Controlled Surface Chemistry

The tip–sample interaction force was first measured on chemically modified SAMs obtained on undeformable substrates (silicon wafers) and was compared with that for as-received silicon wafer. Fig. 3.3 shows that AFM measurements in our conditions are sensitive chemical modification of the layers. When jump-off (or jump-in) contact occurs, we measure the corresponding pull-off (or pull-in) deflection. Pull-off deflection values (Dpull-off) increase in the following order: Dpull-offSiCF3 < Dpull-offSiCH3 < Dpull-offSiasreceived < Dpull-offSiester < Dpull-offSiNH2 < Dpull-offSiOH. Knowing the pull-off deflection, one can easily deduce the adhesion force, if the cantilever spring constant k is known (Fadh = kDpull-off); k has been determined according to the method described in Ref. [5]. The pull-off deflection and thus the adhesion force value increases with the hydrophilic character of the surface. The measured adhesion force depends strongly on the tip–sample contact area, which means it depends on the tip radius in the case of undeformable substrates. Sugawara et al. [6] suggested that the adhesion force is proportional to the tip radius. From the DMT theory [7], which establishes a relationship between the adhesion force F and the thermodynamic work of adhesion W0 (F = 2pRW0; R = tip

Fig. 3.3 Experimental DD curves for functionalized silicon wafers.

39

40


Fig. 3.4 Thermodynamic work of adhesion deduced from the DMT theory, versus surface energy of SAMs grafted on wafers.

radius), we can deduce W0 from experimental adhesion forces. Fig. 3.4 shows that W0 decreases linearly with the surface energy deduced from classical wettability measurements. This linear correlation is obtained after correction for capillary forces observed on SiNH2 and SiOH wafers. In air, a tip–surface capillary bridge is formed for SiNH2 and SiOH surfaces and the pull-off mechanism is different from that for the other surfaces. The capillary force is expressed by Eq. (1), where cw is the surface tension of water, R the tip radius and h the contact angle between water and the tip or water and the wafer. Fcap 2pRcw cos hw=tip cos hw=wafer

1

3.3.3 Force–Distance Measurements on Polymers

A comparison between the DD curves obtained on a silicon substrate and on a PDMS substrate clearly shows the significant contribution of the mechanical properties of the polymer (Fig. 3.5). First, compared with that for silicon wafers, in the tip–polymer separation process the jump-off contact occurs over a large piezo displacement scale and could correspond to a progressive dewetting of the tip by polymer chains during retraction. Secondly, the jump-off amplitude is higher than for silicon wafers. Finally, the loading and unloading slope in the DD representation is much lower than unity in the case of soft polymer systems. The beginning of the indentation is assumed to be at the minimum of the DD curve. Creep experiments [9] have also been performed with the AFM. Considering the experimental contact time (texp < 0.1 s), the creep effect is negligible in our force curve measurements.

3.3.3.1 Force–Indentation Measurements on Polymers Before monitoring DD curves on PDMS, the actuator and the cantilever were thermally stabilized. The laser spot in contact with the tip was positioned in such a way that tangential forces, due to frictional forces, were minimized. We carried out these measurements on the basis of the above prerequisites so that force measurements could be reliable and comparable. Force–indentation (F–I)

3.3 Results and Discussion Fig. 3.5 Comparison between a DD curve obtained on a silicon substrate (top) and a PDMS substrate (bottom).

curves (Fig. 3.6) are deduced from DD curves by assuming that for a given force, the indentation depth is the difference between the experimental deflection value (ds) and the one that would be observed if the material were undeformable (dr deduced from the slope of 1 for undeformable materials). To deduce work of adhesion from our experimental values, it is necessary to have a good estimate of the radius of the probe and thus of the tip–sample contact area. The SEM (Scanning Electron Microscopy) picture reveals that the tip shape can be represented as a cone ended by a sphere. Fig. 3.6 shows that the F–I curve evolves according to two regimes. In our case, the transition between the two modes corresponds to the change in the contact geometry (sphere– plane contact to cone–plane contact). The observed transition shows that our tip radius can be estimated to be 50 ± 5 nm. One has to ask whether the stiffness of the cantilever is suitable. Considering the normal stiffness of the cantilever, it is possible to determine the maximal indentation depth (dmax) by using Eq. (2), where ktip is the constant stiffness of the cantilever, E * is the reduced modulus and R is the tip radius. p 2 ktip 2 E Rdmax Considering that for our substrates the maximum indentation depth is 89 nm (Table 3.3), we have decided to perform constant nano-indentation depth experiments up to 80 nm. Moreover, considering the longitudinal stiffness of the can-

41

42

3 Adhesion at the Nanoscale: an Approach by AFM Fig. 3.6 Experimental F–I curves obtained for the two PDMS having the greatest differences in their Mc values (1.5k and 53k).

Table 3.3 Calculation of the theoretical maximum indentation depth of the tip for our polymer networks.

PDMS

dmax [nm]

1.5k 7k 11k 53k

89 1 098 5 000 26 627

tilever (kx), the tip sliding effect before indentation is assumed to be negligible compared with the total indentation depth. 3.3.3.2 Force–Indentation Curves on Systems of Controlled Surface Chemistry and Controlled Mechanical Properties

All F–I curves obtained on PDMS and grafted PDMS have been performed at an 80 nm indentation depth, as discussed previously. We now consider PDMS of different Young’s moduli, grafted with molecules identical to those for the SAMs on silicon wafers (CH3 and NH2 SAMs). The ratios (FadhPDMSX/FadhSi) (where FadhPDMSX and FadhSi represent the adhesion force measured on PDMS and silicon substrates, respectively, and X

3.3 Results and Discussion Table 3.4 The ratio FPDMSX/FSi for X = CH3, NH2, and as-received substrates, and for three PDMS of differing mechanical properties.

Si-O-Si

CH3 grafts

NH2 grafts

Ratio

FPDMS/FSi

Ratio

FPDMS/FSi

Ratio

FPDMS/FSi

1.5k/Si 11k/Si 53k/Si

1.1 1.8 2.0

1.5kCH3/SiCH3 11kCH3/SiCH3 53kCH3/SiCH3

1.0 1.5 1.9

1.5kNH2/SiNH2 11kNH2/SiNH2 53kNH2/SiNH2

1.2 1.5 2.6

represents the functionality of the grafting) were calculated for all the grafted PDMS and PDMS substrates (Table 3.4) and compared. It appears that for a given PDMS substrate (for example, 1.5k), the ratios are independent of the surface chemistry, whereas for a given grafting, the ratios depend on the mechanical properties of the substrate, with the same values for all the grafts. Thus, for a given substrate, Eq. (3) holds (k is a constant). FadhPDMSX Fadhsix k

3

However, for an undeformable substrate, the DMT theory [7] gives Eq. (4), and hence Eq. (5). Fadhsix 2pRW0

4

FadhPDMSX 2pRW0 k

5

Introducing a dimensional constant (k' = (2pR)–1), this relationship becomes Eq. (6) or (7), where G is the separation energy and f (Mc, v, T) is a viscoelastic dissipation function which depends on the network molecular structure (Mc), temperature T and separation rate v. GadhPDMSX W0 k 2pR k0

6

GadhPDMSX W0 f MC; v; T

7

or

This relationship clearly distinguishes between the mechanical and chemical contributions in a force–ndentation measurement with an AFM. According to Table 3.5, f (Mc, v, T) is determined for each substrate for a given rate and a given temperature and is independent of the surface chemistry. From the theoretical point of view, f (Mc, v, T) cannot be lower than 1 (which corresponds to a zero separation rate v) and should increase while the Young’s modulus decreases (which means energy dissipation in the bulk is higher when the network is softer). The values shown in Table 3.5 are coherent with these as-

43

44

3 Adhesion at the Nanoscale: an Approach by AFM Table 3.5 Average viscoelastic dissipation function deduced from calculated thermodynamic work of adhesion.

PDMS

f (Mc, v, T)

W0CH3 (mJ/m2)

W0 NH2 (mJ/m–2)

1.5k 11k 53k

1.1 1.6 2.2

35 36 32

107 105 106

Fig. 3.7 Dissipation function versus mass between crosslinks.

sumptions. Moreover, even if the dependence of f on the separation rate v is not dominant in the range of the rates available with the AFM, it is obvious (Fig. 3.6) that dissipation occurs for the 53k sample and that this dissipation is included in the function f. Values of the dissipative function f (at room temperature and separation rate of 6 lm s–1) and thermodynamic work of adhesion W0 are listed in Table 3.5. The function f is constant for a given grafting and whatever the substrate Mc value. Values of W0 are also in agreement with values quoted in the literature. Indeed, usual work of adhesion values lie between 40 and 70 mJ m–2 for organic–organic contacts, such as between two silanated silicas [10], and between 40 and 145 mJ m–2 for a contact between a raw material such as silica and a silanated silica. Finally we have reported in Fig. 3.7 the evolution of the dissipative function f versus Mc, the mass between crosslinks. A linear relationship is obtained. Therefore a modified expression for G, the separation energy, is proposed (Eq. (8), where f '(v, T) represents a viscoelastic dissipation function which depends only on the temperature T and separation rate v). GadhPDMSX W0 1 Mc f 0 v; T

8

References

3.4 Conclusion

The model systems studied have allowed us to express the mechanical and chemical surface contributions in a force curve measurement and to establish a relationship at the nanoscale which is quite similar to the relationship of Gent and Schultz [11]. Then a new relationship has been proposed to determine the thermodynamic surface properties of viscoelastic materials on the basis of AFM experiments.

References 1 Mougin K., Haidara H., Castelein G. Col-

2 3 4 5

6

loids and Surfaces: Physicochemical and Engineering Aspects 2001,193, 231–237. Chaudhury M.K. , Whitesides G.M. Science 1992, 255, 1230–1232. Vonna L. Ph. D. Thesis, Université de Haute Alsace, Mulhouse, France 1999. Flory P. J. Chem.Rev. 1944, 35, 51. Torii A., Sasaki M., Hane K., Okuma S. Measurement Science and Technology 1996, 7, 179–184. Sugawara Y., Ohta M., Konishi T., Morita S., Suzuki M., Enomoto Y. Wear 1993, 168, 13–16.

7 Derjaguin B. V., Muller V. M., Toporov

Y. P. J.Colloid Interface Sci. 1975, 53, 314. 8 Basire C. Ph D. Thesis, ESPCI–Univer-

sité Paris VI, Paris, France 1997. 9 Johnson K. L., Kendall K., Roberts A. D.

Proc. R. Soc. 1971, A324, 301. 10 Papirer E., Balard H., Sidqi M. J. Colloid

Interface Sci. 1993, 159, 238–242. 11 Gent A., Shultz J. J.Adhesion 1972, 3,

281.

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47

4 Organization of PCL-b-PMMA Diblock Thin Films: Relationship to the Adsorption Substrate Chemistry T. Elzein, M. Brogly, and J. Schultz

Abstract

By combining polarization-modulation infrared reflection–absorption spectroscopy (PM-IRRAS) and atomic force microscopy (AFM) performed in tapping mode at different tapping forces we were able to define the model of organization and chain folding in thin films of polycaprolactone-b-poly(methyl methacrylate) diblock copolymers (PCL-b-PMMA) spin-coated on gold-coated substrates. Attention was also given to understanding the selective adsorption of the diblocks when the chemistry of the adsorption surface was varied. Our results showed a perpendicular orientation of the PCL block with respect to the adsorption substrate, and that the nucleation takes place on the gold interface during PCL crystallization with rejection of PMMA chains outside the crystallite. Changes in the adsorption model were detected when these systems were adsorbed onto OH-functionalized gold substrates.

4.1 Introduction

The design of new adhesives and/or the creation of model surfaces for further application in adhesion could be improved by the synthesis of new polymers having complex structures with controlled architecture, such as block copolymers. The interest of the latter is widely discussed in the literature – particularly their ability to organize into nanometer-sized domains, thus making them suitable for surface patterning [1–7]. Many parameters influence the characteristic dimensions of the resulting morphology of block copolymers [7–11]: the molecular weight, the volume fraction of the components, the interaction between the blocks, and finally the molecular architecture of the block copolymer. Polycaprolactone-b-poly(methyl methacrylate) (PCL-b-PMMA) diblocks with various compositions of PCL and PMMA fractions are the subject of this study: the PCL is a semi-crystalline polymer known for its biocompatibility, its biodeAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

48

4 Organization of PCL-b-PMMA Diblock Thin Films

gradability, and its use in biomedical domains [12–16]. PMMA, an amorphous polymer (glassy at ambient temperature) is widely used in various fields of application. During this work, we analyzed the organization, the morphology, and the preferential adsorption of these copolymers and how these properties would be influenced by the interfacial interactions that could be established between the polymer and the chemical functionalities grafted onto the adsorption substrate. Our results were mainly based on two surface techniques: polarization-modulation infrared reflection–absorption spectroscopy (PM-IRRAS) and atomic force microscopy (AFM).

4.2 Materials and Methods 4.2.1 PCL-b-PMMA Diblocks

Two PCL-b-PMMA diblocks with different compositions were the subject of this study (see Table 4.1) [17]: · Diblock 1 with a ratio of monomers (m/n) of 48 : 75 · Diblock 2 with m/n = 50 : 30. The general formula of these diblocks is 1.

Thin films of PCL-b-PMMA were spin-coated from 2 g L–1 solutions in toluene on gold-coated glass slides and then on OH-functionalized gold substrates (by grafting of OH-terminated thiols).

Table 4.1 Composition of the PCL-b-PMMA diblocks.

Diblock

m (PCL monomers)

n (PMMA monomers)

PCL weight percentage vPCL

1 2

48 50

75 30

41.4 61.7


Gold substrates were prepared by using glass slides cleaned with ethanol, treated with argon plasma for 3 min at a power of 80 W, and heated at 50 8C for 30 min in a Piranha solution. A gold (99.99% from Balzers) layer (90 nm) was then coated under vacuum (10–7 mbar) by the PVD technique after immersion (12 h) in 3-mercaptopropyltriethoxysilane solution (in toluene). The resulting roughness of the gold layer was in the 1–2 nm range. 3-Mercaptopropyltriethoxysilane (95%) from ABCR was used as a coupling agent between glass slides and gold coatings. For OH grafting, gold substrates were cleaned by UV radiation before being immersed for about 12 h in 11-mercapto-1-undecanol (97% from Sigma-Aldrich) solution (in 99.99% ethanol). A typical concentration of 3 mM was chosen for the grafting solution. The thickness, as obtained from ellipsometric measurements, ranged between 10 and 14 nm for Diblock 1 and Diblock 2 thin films on gold substrates and between 5 and 8 nm on OH substrates. 4.2.2 Infrared Spectroscopy 4.2.2.1 Transmission Powder samples of diblocks were analyzed by infrared spectroscopy with the transmission device. Measurements were performed with an IFS48 Bruker spectrometer and the number of scans was fixed at 100 with a resolution of 4 cm–1.

4.2.2.2 Polarization-Modulation Infrared Reflection–Absorption Spectroscopy (PM-IRRAS) Diblock thin films were analyzed by PM-IRRAS. Measurements were taken with an IF66S Bruker spectrometer; spectra were recorded with an MCT detector, under experimental conditions of 1000 scans, a resolution of 1 cm–1, 858 as the beam angle of incidence and a 74 kHz modulation frequency. A ZnSe photoelastic modulator provided by Hinds Instruments was used. The basic principle of the PM-IRRAS method (Fig. 4.1) is to combine the FTIRRAS experimental conditions with a fast modulation of the polarization state of the incident electric field (ideally between p and s linear states) and to extract from the detected intensity (using electronic filtering and demodulation) the two signals (Rp–Rs) and (Rp + Rs) in order finally to compute the differential reflectivity spectrum DR/R [18]. The basis and principles of this technique are detailed elsewhere [18–24]. However we note the following points: · The selection rules induced by the reflection of a p polarized infrared beam under grazing angles on metallic surfaces imply that when a dipole moment is oriented normal to the surface, its signal is intensified in the PM-IRRAS spectrum.

49

50


Fig. 4.1 PM-IRRAS setup.

· The double modulation of the incident beam and the mathematical treatment of the detected signal allow a differential reflectivity (DR/R) surface signal to be obtained in a single step and with the whole dynamic range of the detection. · An anisotropic phase has a highly intensified signal in the PM-IRRAS spectrum; in fact, the isotropic phase makes a contribution to the PM-IRRAS global spectrum but its signal is very weak compared with the anisotropic phase. · And finally, due to the main points described above, PM-IRRAS offers the possibility of accessing quantitative calculation of the orientation angles at the surface if the qualitative orientation model is well defined. 4.2.3 Atomic Force Microscopy (AFM)

Measurements were performed with a Nanoscope IIIa/Dimension 3000 combination (Digital Instruments) in the tapping mode and under ambient conditions. The electronic extender module allows simultaneous phase detection and height imaging. Experimental details are given in the Results and Discussion section.

4.3 Results and Discussion 4.3.1 PCL-b-PMMA Bulk Characterization

The degree of crystallinity of the diblock, Xcdiblock, and that related to the PCL entity, XcPCL (XcPCL = Xcdiblock/vPCL) were calculated on the basis of DSC analysis. XcPCL was found to be 55% and 72% respectively for Diblock 1 and Diblock 2, while it was 77% in the case of pure polycaprolactone.


It is obvious that the presence of the PMMA block (amorphous) leads to a decrease in the degree of crystallinity of the PCL block; this is more evident in Diblock 1 having 58.6% (w/w) of PMMA, than in Diblock 2 having only 38.3% of PMMA. However, this effect is expected due to the steric overlapping of the PMMA block during the crystallization of the PCL chains: the amorphous chains of PMMA influence the chain folding in the PCL block, thus leading to a reduction in crystallinity of the latter. These diblocks, as well as pure PCL, were also analyzed by infrared spectroscopy in the transmission mode in order to identify all vibrators, their assignments, and their position [25]. Moreover, these spectra are considered as the isotropic response of these materials. This fact is of major importance when discussing, later in this paper, the orientation of the chains in thin films of the diblocks. 4.3.2 PCL-b-PMMA Thin Films on OH-Functionalized Gold Substrates

Experimentally, to access the organization of chains in thin films, the initial step is comparison of the spectra of the bulk and thin films since, according to the selection rules, the changes in band intensity are related to the preferential orientation of the chains. In that way, we compared the PM-IRRAS spectrum of bulk Diblock 2 with that of its thin film (Fig. 4.2). In Fig. 4.2 one can observe the important decrease in intensity of the band around 1150 cm–1 in the thin film-spectrum. This band is attributed to the amorphous phase and the majority of it is a PMMA contribution [25]. Moreover, the bands related to the CH3 stretching modes in the 2500–3000 cm–1 region

Fig. 4.2 Comparison of bulk and thin-film PM-IRRAS spectra of Diblock 2.

51

52


are no more clearly distinguished in the thin film spectrum. As a consequence we can suppose, for instance, that the PMMA chains are flattened on the gold surface and are adopting a conformation close to the isotropic one. But how should the entire diblock thin-film spectrum be interpreted? To answer this question, we compared the PM-IRRAS spectra of the pure PCL and the Diblock 2 thin films (Fig. 4.3). In previous work dedicated to polycaprolactone [26], we developed an adsorption model of pure PCL in which the main chain axis CC and the CCC plane of the PCL were almost perpendicular to the substrate surface and that for different chain lengths. This preferential orientation could probably be imposed during thin-film crystallization at the gold surface since that PCL seems to remain crystalline even in the confined geometry (as evidenced by appearance of the crystalline band in the thin film spectrum). Thus in the case of the diblock thin-film spectrum being identical to that of the pure PCL thin film, we suggest the same preferential orientation of the PCL crystalline chains in the diblock thin film. Even if many bands are common to both PCL and PMMA, the chain flattening of the PMMA block would induce disappearance of the PMMA contribution in the thin film spectrum (according to PM-IRRAS selection rules), and this could be justified by the amorphous nature of the PMMA (in the absence of important forces such as crystallization to promote a preferential orientation). Thus, the resulting spectrum of the diblock thin film is considered, in a first approximation, as the contribution from the PCL crystalline phase. As a consequence, we suggest an adsorption model of the PCL-b-PMMA copolymers on the OH grafted gold surface in which the PCL crystalline chains are oriented almost in the normal direction (with respect to the gold substrate surface), while no preferential orientation is observed for the PMMA block.

Fig. 4.3 Comparison of pure PCL and Diblock 2 thin-films PM-IRRAS spectra.


However, a complete description of this model could not be given without an AFM morphological analysis of thin films of the diblocks. AFM experiments are carried out by AFM using a Dimension 3000 microscope coupled to a Nanoscope IIIa electronic controller (Digital Instruments, Veeco-FEI Co., USA). All experiments were performed in tapping mode. The AFM was equipped with the phase extender electronic modulus, making it possible to record the phase shift variations between the instantaneous oscillation of the tip and the oscillation applied to the cantilever in tapping mode. In our case, this phase shift depended strongly on the local moduli between the different components of the material and reflected the surface structure of the thin films [27–34]. It is also important to mention that the measurements were also performed at different amplitudes in order to probe surface rigidity versus indentation on the basis of the phase image. Nevertheless, amplitude reduction never exceeded 15%, corresponding to relatively low levels of force exerted on the surface (“light tapping”) [30–35]. In this mode, topographic images are supposed not to be perturbed by mechanical effects. Then differences on phase images can be interpreted properly. It should be noted that semi-crystalline polymers are well suited for the use of the phase mode as surface viscoelastic properties between crystalline and amorphous regions are large. Fig. 4.4 represents a comparison of topographic images (10 lm ´ 10 lm) for two samples of PCL-b-PMMA thin films on OH-gold substrates (Diblock 1 and Diblock 2). Both images show characteristic fingerlike patterns. The number of fingerlike patterns increases when the length of the PCL block increases relative to the PMMA one. It then appears obvious that the fingers correspond to the organiza-

Fig. 4.4 AFM topographic images of Diblock 1 and Diblock 2 thin films on OH substrates (10 lm ´ 10 lm). The z range is 20 nm.

53

54


tion in a crystalline phase of the PCL blocks. The difference in height between the apparently amorphous phase and the crystalline one is 6 nm, as confirmed by measurement of the cross-section. However, a complete description of block distribution and organization could not be given from a simple observation of morphological features. In order to identify the difference in the surface rigidity of the PCL-b-PMMA thin films we have performed tapping-mode AFM measurements at different tapping forces on the Diblock 2 sample. Three parameters must be considered in estimating the tapping force: the free amplitude A0; the set point ratio rsp; and the cantilever spring constant k. The free amplitude is the oscillation amplitude of the cantilever when there is no interaction with the surface of the sample. The set point ratio equals Asp/A0 where Asp, the set point amplitude, is the amplitude of the cantilever when it has been reduced by contact with the sample. In the case of a polymer surface that is prone to surface indentation, Asp is equivalent to the tip–surface distance (dsp) plus the indentation depth of the tip into the sample (zind). Changing the rsp value has been found to lead to great differences in the phase images obtained [36, 37]. To reduce the force applied to the surface of the sample, rsp should be close to unity. As rsp is lowered it is possible to image subsurface features [36–38]. In that way Fig. 4.5 shows a characteristic PCL-b-PMMA (Diblock 2) fingerlike pattern imaged at different rsp values: 0.96, 0.86, and 0.83. As rsp decreases, the amplitude of the tip oscillation increases and topographic images are rather similar, but great differences are observed on phase shift images: the difference in mechanical properties of the different regions emerges clearly. On the color scale, the dark zones correspond to soft regions and the bright zones to hard ones. As fingerlike patterns are attributed to crystalline organization of the PCL blocks, it seems reasonable to assume that the phase between the fingers is the amorphous PMMA one. In Fig. 4.5(a), at small tapping force, the region between fingers is rather dark (i.e., soft) and as the tapping force increases (Figs. 4 b and c) brighter zones emerge delimiting regions with different mechanical properties that correspond to the finger morphology. This observation allows us to confirm that regions between the fingerlike patterns correspond to the amorphous PMMA blocks and thus when the tapping force increases this layer is indented and the tip becomes sensitive to the gold substrate (bright), which has greater rigidity than the PMMA one. On the other hand, at high and low rsp values fingers appear to be unchanged due to the free surface of the PCL crystallites (there is no PMMA layer covering the PCL fingers). According to this scheme, crystallization of PCL is initiated at the substrate interface. We can also estimate to a first approximation the height of the PMMA layer on the basis of the phase images at different set point ratios. A reduction of set point ratio from 0.96 to 0.83 corresponds to a reduction of probing amplitude of 3 nm (the free amplitude is initially set to 24 nm). This result means that the thickness of the PMMA layer is 3 nm (Fig. 4.6 b), since no changes are observed when working with rsp lower than 0.83, and it means that the PMMA


Fig. 4.5 AFM images of Diblock 2 thin films on OH substrates. (a) rsp = 0.96; (b) rsp = 0.86; (c) rsp = 0.83.

layer is almost fully indented. On the other hand, the analysis of the cross-section in AFM images showed that the difference between the fingers and the region between them is about 6 nm, which means that the PCL crystallites are about 9 nm thick. 4.3.3 PCL-b-PMMA Thin Films on Gold Substrates

The same analyses were conducted on the diblock thin films adsorbed on inert gold substrates [25]. No great differences were revealed in chain orientation when comparing to films on OH substrates. Meanwhile, our AFM analysis showed almost the same general morphology, but when working under different tapping forces [25] we noticed that fingers appear dark at small tapping

55

56


Fig. 4.6 Adsorption model of PCL-b-PMMA: (a) on a gold surface; (b) on an OH-grafted gold substrate.

forces and as rsp decreases they become brighter, leading us to suspect that a thin amorphous PMMA film could also be present on the surface of the crystalline PCL fingers and that at low rsp the PMMA layer is indented and the tip becomes sensitive to the PCL crystalline phase (brighter), which is stiffer than PMMA and softer than the gold substrate (Fig. 4.6 a). The difference between these two models could be due to the acid–base interactions established between OH grafts and PMMA chains: during evaporation of the solvent (spin-coating) PMMA blocks will be attracted to the OH grafts and thus will be more flattened on the functionalized gold interface, while the PCL chains implied in the crystalline phase are less mobile and are not able to establish such interactions.

4.4 Conclusion

This work was dedicated to a study of the organization of PCL-b-PMMA diblock copolymers adsorbed on OH-functionalized gold-coated substrates in order to analyze the effects induced by confinement and interfacial interaction compared with inert substrates. The access to these parameters was based on two surface techniques that are highly suitable for thin-film analyses – PM-IRRAS and AFM. Our results enabled us to prove that the orientation of PCL crystalline chains is not affected by the presence of the PMMA block, by comparison with pure

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Spectrosc. 1991, 45, 380. 19 T. Elzein, M. Brogly, G. Castelein,

J. Schultz, J.Polym. Sci. B 2002, 40, 1464. 20 T. Elzein, M. Brogly, J. Schultz, Polymer

2002, 34, 4811. 21 T. Elzein, M. Brogly, J. Schultz, J. Surf.

Interf. Anal. 2003, 35, 231. 22 T. Elzein, M. Brogly, J. Schultz, Polymer

2003, 44, 3649. 23 T. Elzein, M. Brogly, J. Schultz, J. Surf.

Interf. Anal. 2003, 35, 633. 24 T. Buffeteau, B. Desbat, E. Péré, J. M.

Turlet, Mikrochim. Acta 1997, 14, 627. 25 T. Elzein, H. Awada, M. Nasser-Eddine,

26

27 28

29

30 31

32

C. Delaite, M. Brogly, Thin Solid Films 2005, 483, 388. T. Elzein, M. Nasser-Eddine, C. Delaite, S. Bistac, Ph. Dumas, J. Coll. Interf. Sci. 2004, 273, 381. T. Buffeteau, B. Desbat, E. Péré, J. M. Turlet, Mikrochim. Acta 1997, 14, 627. G. Bar, R. Brandsch, M. Bruch, L. Delineau, M. H. Whangbo, Surf. Sci. 2000, 444, L11. G. Bar, L. Delineau, R. Brandsch, M. Bruch, M. H. Whangbo, Appl. Phys. Lett. 1999, 75, 4198. S. N. Magonov, V. Elings, M. H. Whangbo, Surf. Sci. 1997, 375, L385. D. Sarid, T. G. Ruskill, R. K. Workman, D. Chen, J. Vac. Sci. Technol. B 1996, 14, 864. D. Krüger, B. Anczykowski, H. Fuchs, Ann. Physik 1997, 6, 341.

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4 Organization of PCL-b-PMMA Diblock Thin Films 33 S. N. Magonov, V. Elings, V.S. Papkov,

Polymer 1997, 38, 297. 34 G. Bar, Y. Thomann, R. Brandsch, H. J. Cantow, M. H. Whangbo, Langmuir 1997, 13, 3807. 35 J. Mallégol, O. Dupont, J. L. Keddie, Langmuir 2001, 17, 7022.

36 A. Knoll, R. Magerle, G. Krausch, Macro-

molecules 2001, 34, 4159. 37 A. Buzin, Y. K. Godovsky, N. N. Makaro-

va, J. Fang, X. Wang, C. M. Knobler, J. Phys. Chem. B 1999, 103, 11 372. 38 J. Feng, L. T. Weng, C. M. Chang, J. Xhie, L. Li, Polymer 2001, 42, 2259.

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5 Adhesion and Friction Properties of Elastomers at Macroscopic and Nanoscopic Scales S. Bistac and A. Galliano

Abstract

Friction and adhesion of elastomers are quantified at two different scales: at the macroscopic scale, using a tack test and a pin on disk tribometer, and at the nanoscopic scale, using atomic force microscopy (AFM). The objective is firstly to verify if friction and adhesion behaviors are comparable at both scales, and secondly to define molecular mechanisms that are able to explain the measured properties. The influences of structural parameters (degree of crosslinking and presence of free chains) and experimental factors (separation speed, normal force, and initial contact time) are analyzed. The polymers are crosslinked polydimethylsiloxane (PDMS) and the substrate (for the macroscopic tests) is a smooth glass plate. Experimental results underline the major role of molecular parameters such as degree of crosslinking and length and content of free and pendant chains. Conceptual schemes are proposed to describe the complex interfacial molecular mechanisms.

5.1 Introduction

Dissipation phenomena generally occur during measurement of the adherence of polymer materials, leading to an adherence energy function of both the number and nature of interfacial interactions (adhesion) and dissipative properties, mainly due to viscoelastic behavior [1–5]. Friction properties of polymers are also governed by interfacial interactions and dissipation mechanisms. Common phenomena (interfacial interaction and dissipation) therefore control adherence and friction behaviors. However, the relationship between the two phenomena is still vague or undefined. The first objective of this experimental work is then to compare adherence and friction of polydimethylsiloxane (PDMS) networks in order to establish relationships between these two properties. Adherence and friction are generally measured at the macroscopic scale. However, the development of atomic force microscopy (AFM) for the measureAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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5 Adhesion and Friction Properties of Elastomers at Macroscopic and Nanoscopic Scales

ment of mechanical response at a nanoscale allows us to determine adherence and friction between an AFM tip and the PDMS surface. The second objective of this study is then to compare adherence and friction behaviors of PDMS at both scales, “macro” and “nano”, in order to verify if the involved mechanisms are similar and if the nano-scale approach allows a better understanding of what is happening at a macroscopic scale. In practice, tack experiments, translation tribometry, and atomic force microscopy (AFM) are used to quantify adhesion and friction at the macro- and nanoscales. In order to record dissipation phenomena, the influence of structural parameters (degree of crosslinking and presence of free chains) and experimental factors (friction speed, normal force) is analyzed.


Two PDMS with different initial molecular weights (Mw = 6000 g mol–1 (called PDMS 6) and Mw = 17 200 g mol–1 (PDMS 17) were used. They were vinyl-terminated and crosslinked with tetrakis(dimethylsiloxy)silane. Increasing the initial molecular weight induced two major consequences: a lower crosslinking density and also a greater quantity of free chains and pendant chains (linked to the network by only one extremity). Free chains could be extracted by immersion in a good solvent. Networks have also been studied after extraction of free chains (samples PDMS 6' and PDMS 17'). The glass temperature of PDMS networks is –122 8C. All experiments were performed in ambient air, at room temperature (20 8C). The substrate used for tack experiments and macroscopic friction was a smooth glass plate, cleaned with ethanol and dried before use (Fig. 5.1). During the tack experiment, the polymer hemisphere (diameter = 16 mm) was brought into contact (at a fixed approach speed of 10 mm min–1) with the glass substrate, during a given contact time and under a controlled normal load. Both materials were then separated at a given speed, and the separation force was measured. The apparatus simultaneously determined the force and the apparent contact area between the substrate and the polymer, using a video camera placed under the transparent glass substrate. Friction properties of PDMS hemispheres were measured using a translation tribometer. The PDMS hemisphere was brought into contact with the glass plate, under a given normal load (Fig. 5.2). The glass plate was then moved in translation at a controlled speed and the tangential force (= friction force) was measured. The evolution of the contact area was also followed with a video camera. The adherence energy, expressed in J m–2, was calculated by dividing the integral of Fds versus distance by the contact area corresponding to the maximum force measured during separation. Nanoscale adhesion and friction of PDMS films were quantified by AFM using a Dimension 3000 microscope (Digital Instruments) in contact mode with


sensor

Fig. 5.1 The tack apparatus.

a commercial silicon tip on a 100 lm triangular cantilever (spring constant = 0.58 N m–1). Force–distance curves obtained during a loading–unloading experiment allowed the determination of nano-adhesion, which is directly proportional to the maximum cantilever deflection during the tip/PDMS separation, as illustrated in Fig. 5.3. The nano-friction force, measured in torsion mode, is directly proportional to the TMR value (trace minus retrace, in volts), which is given as the difference between the lateral forces scanning left-to-right and right-to-left. Absolute adhesion and friction forces can be obtained with calibration methods but such techniques

sensor

Fig. 5.2 The translation tribometer.

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Fig. 5.3 Force–distance curve obtained by AFM (nano-adhesion). Fadhesion = KD, where Fadhesion = adhesion force; D = maximum deflection; K = calibration constant.

are not necessary for this comparative study. Cantilever deflection in nanometers (or adhesion force) and TMR values in volts (or friction force) have then been compared for both polymers, before and after extraction of free chains.


Wettability measurements indicated a similar surface energy for PDMS 6 and 17, before and after extraction, close to 27 mJ m–2, with a dispersive component equal to 27 mJ m–2 and a nondispersive component equal to zero. The identical surface energy and chemical composition showed that the ability of both polymers to interact with the substrate was then theoretically identical. 5.3.1 Adherence Energy

Table 5.1 presents the adherence energy of PDMS samples, before and after extraction of free chains, measured by tack. Surface analysis performed by the IRRAS (infrared reflection–absorption spectroscopy) technique and AFM on the glass substrate after separation indicated that the failure occurred inside the polymer, very close to the interface with the glass substrate. AFM and infrared spectroscopy have indeed evidenced the presence of a thin residual polymer layer (thickness close to 10–20 nm) on the glass plate after the tack test. The thickness of this layer was increased for higher molecular weight. Despite the presence of this very thin residual layer after separation, we will consider, in a first approximation, the measured energy as the adherence energy, representative of the interfacial strength. Due to this transfer layer, a new glass plate and PDMS hemisphere have been used for each experiment.

5.3 Results and Discussion Table 5.1 Adherence energy E [J m–2] of PDMS 6 and 17 (before extraction) and PDMS 6' and 17' (after extraction) measured for a normal force of 1 N and for an initial contact time of 300 s and two separation speeds V (DE = 0.02 J m–2).

Separation speed V [mm min–1] 1 PDMS 6 PDMS 17 PDMS 6' PDMS 17'

0.51 1.79 1.21 1.70

100 1.04 3.69 2.04 7.00

Tack results indicated a higher adherence energy for PDMS 17 than for PDMS 6. Separation speed also greatly affected the adherence value, with an increase of the separation energy with the separation speed [6]. Even if the chemical composition and surface energy of polymers were identical, the PDMS samples studied differed greatly in their chain lengths, degree of crosslinking, and content of free and pendant chains. The higher mobility of the chains of PDMS 17 (more free and pendant chains) will induce a better contact with the substrate, at a molecular scale. This greater chain adsorption of PDMS 17 onto the substrate will increase the number of interactions per area with the glass substrate. However, the adherence energy measured during the separation is also a function of the polymer dissipative properties. During separation, the chains are extended before contact failure. This motion, disentanglement, and breaking of chains is more dissipative for PDMS 17, due to the longer chains. Both phenomena, i.e., a greater density of interfacial interactions and more dissipative chain motion, explain the higher adherence of PDMS 17. The influence of speed is significant, as shown by the increase of adherence energy. It is explained by the viscous behavior of large-scale chain motion during separation. This large-scale motion, which affects adsorbed chains (free and pendant) preferentially, is rate-sensitive, due this viscous behavior. Movements and pull-out of longer chains are more difficult and more sensitive to the separation speed. After extraction of free chains, the adherence energy is higher for both grades of PDMS. Firstly, network chains no longer compete with free chains (which are now absent) for adsorption sites on the glass substrate. Secondly, pendant chains exhibit a higher mobility than crosslinked chains (possible favored adsorption onto the substrate) but are also chemically bonded to the network, allowing more stress to be transferred than free chains do. Therefore these pendant chains play a major role in adhesion behavior [7]. A thin residual PDMS layer is detected by AFM and IRRAS after friction on the glass substrate [8].

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5.3.2 Macroscale Friction

To determine the macro-friction properties, the tangential or friction force was measured by a tribometer, and the contact area evolution was recorded by video camera. Friction stress was calculated by dividing each force value by the corresponding contact area. Results of the macro-friction experiments are presented in Table 5.2. The first and main result is that a higher friction stress is measured for PDMS 6 than for PDMS 17. Tack results have shown that PDMS 17 exhibits the greater adherence. This higher adherence does not induce a higher friction, however. The presence of a permanent contact during friction (induced by the applied normal load), can explain these results. This elastic contact acts like a “forced wetting” and therefore compensates for the low adhesive contact of PDMS 6. Moreover, for PDMS 6, the lower content of free and pendant chains induces a more direct and efficient contact with the glass substrate. The crosslinked network will then be directly constrained during friction, thus increasing the bulk dissipation. For PDMS 17, chain orientation phenomena (especially for free and pendant chains) induced by the interfacial shear stress during friction are important due to the greater length of the chains [8]. This orientation and alignment of chains at the interface can act like a self-lubrication layer, hindering the direct contact between the glass substrate and the bulk network, and therefore decreasing the bulk dissipation and the friction resistance. The effect of extraction on macro-friction is globally slight, with almost identical values for PDMS 6 and 17 before and after extraction. Friction speed has a small effect on the maximum friction stress. This low dependence on speed can be explained by the very low value of the glass transition temperature: PDMS networks are used at ambient temperature, on their rubbery plateau. However, for PDMS 6, a slight decrease in friction stress could be explained by a pseudoplastic behavior of the confined interfacial layer (a shear thinning effect).

Table 5.2 Friction stress r [N mm–2] for different friction speeds (Dr = 0.01 N mm–2).

Friction speed [mm s–1] 0.4 PDMS 6 (before extraction) PDMS 17 (before extraction) PDMS 6' (after extraction) PDMS 17' (after extraction)

0.22 0.07 0.22 0.07

2 0.19 0.07 0.21 0.07


To summarize, macroscale adhesion results have shown a higher adherence for PDMS 17 and a large effect of extraction and separation speed. Macroscale friction results have indicated a higher friction resistance for PDMS 6 than for PDMS 17, and a slight effect of extraction and friction speed. 5.3.3 Nanoscale Friction and Adhesion

Nanoscale measurements have been performed by AFM in contact mode. For the nano-adhesion measurement, force–distance curves have been obtained (Fig. 5.3). The maximum cantilever deflection D during separation is directly proportional to the adhesion force. For PDMS 6, the deflection D was quite constant before and after extraction and was close to 70 nm (with DD = 5 nm). The deflection value did not depend on the contact force for PDMS 6. Measurements were not possible with PDMS 17 before extraction because the AFM tip could not be separated from the PDMS surface within the measurable cantilever deflection range. Hence the adhesion value could not be determined, but it has to be considered as important. After extraction of free chains (PDMS 17'), measurements were possible and the deflection value D obtained was close to 500 nm (with DD = 30 nm). Hence, in correlation with the macro-adhesion results, the nano-adhesion of PDMS 17 was much greater than that of PDMS 6. However, the nano-adhesion of PDMS 17 decreased after extraction. This result is the opposite of that for macro-adhesion, where an increase in adherence was observed after extraction. Important adsorption phenomena of the numerous and long free chains of PDMS 17 on the AFM tip could explain the higher nano-adhesion observed before extraction. The mobility of these free chains (greater than of pendant chains) allows a better adsorption on the tip, avoiding the separation (with the same experimental device, i.e., cantilever stiffness). For nano-friction determination, the TMR was measured; it is directly proportional to the friction force. The influences of normal load (deflection set point, in volts) and friction speed (tip velocity, obtained by varying the scan frequency) were analyzed. Nano-friction measurements were impossible for PDMS 17 (before extraction) because the AFM tip was “trapped” in the surface layer. However, friction between the tip and the surface of PDMS 17 can be considered as very strong, even if it has not been quantified here. AFM measurements were possible for PDMS 6 (before and after extraction of free chains) and PDMS 17' (after extraction of free chains). Figs. 5.4 and 5.5 show the evolution of the TMR as a function of the applied load for different friction speeds on PDMS 6 and 6'. The nano-friction coefficient is proportional to the TMR divided by the applied load. Straight lines are observed for the TMR in the given load range for both samples, indicating an almost constant friction coefficient whatever the applied load. The effects of extraction and speed are strong, as illustrated in Table 5.3, which represents the

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Fig. 5.4 Evolution of the TMR (nano-friction force) as a function of applied load for the different friction speeds indicated, measured for PDMS 6 (before extraction).

Fig. 5.5 Evolution of the TMR (nano-friction force) as a function of applied load for the different friction speeds indicated, measured for PDMS 6' (after extraction).

Table 5.3 Friction force F (in terms of TMR value [V], which is equal to the difference between lateral forces scanning left-to-right and right-to-left) of PDMS 6 and 6' for various speeds and applied loads measured by AFM (DF = 0.01 V).

Speed [lm s–1]

Applied load [V] –0.5

0

1

PDMS 6 (before extraction)

5 50

0.11 0.28

0.13 0.33

0.16 0.39

PDMS 6' (after extraction)

5 50

0.06 0.21

0.09 0.28

0.13 0.40


friction “force” by the TMR, measured for PDMS 6 and 6'. A higher TMR corresponds to a greater friction force. Surprising results, compared with macroscopic friction, were obtained for PDMS 6, with a higher nanoscale friction before extraction of free chains. The influence of extraction was also more pronounced for lower normal loads. Macroscopic tests previously showed a slight increase in friction after extraction. The influence of speed is obvious (contrary to the macroscopic results), with an increase in friction with speed for both grades of PDMS. This effect is similar to the influence of speed observed in macroscale adhesion, for which a higher separation speed induces an increase in the tack energy. Chain motions during nano-friction and pull-out mechanisms, which are speed-dependent, could explain this behavior. Nano-friction measurements were impossible on PDMS 17 before extraction, indicating a high friction resistance. After extraction, the nano-friction could be measured for PDMS 17'; the results are illustrated in Fig. 5.6. The first important result is that nano-friction decreased after extraction – the TMR value could then be measured. This effect can also be explained by the high adsorption of long and numerous free chains on the AFM tip, preventing the sliding of the tip. The straight lines in Fig. 5.6 indicate a constant nano-friction coefficient and the absence of influence of normal load. This figure also indicates that the influence of friction speed was less pronounced than for PDMS 6. Table 5.4, which reports friction values (TMR values) for PDMS 17', shows a slight effect of speed. Comparing the nano-friction of PDMS 6 and 17, that for PDMS 17 (TMR > 0.8) was much higher than for PDMS 6 (TMR < 0.4). This comparison, in contradiction to the macro-friction results, involves the effect of the number and length of mobile chains (free and pendant). The longer and more numerous free and pendant chains present in PDMS 17 will induce a great adsorption

Fig. 5.6 Evolution of the TMR (nano-friction force) as a function of applied load for the different friction speeds indicated, measured for PDMS 17' (after extraction).

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5 Adhesion and Friction Properties of Elastomers at Macroscopic and Nanoscopic Scales Table 5.4 Friction force F (in terms of TMR [V]) of PDMS 17' after extraction of free chains, measured by AFM (DF = 0.1 V)

Speed [lm s–1]

20 100

Applied load [V] –0.5

0

1

1.0 0.8

1.4 0.9

1.8 1.9

on the AFM tip, as in nano-adhesion. Nanoscale behaviors are then governed by adsorption phenomena on the tip, favored by chain mobility [9]. To sum up, nanoscale behavior can be totally different from macroscopic properties. Friction and adhesion properties are very sensitive to molecular parameters such as the degree of crosslinking (or chain length) and presence of free chains. When the degree of crosslinking decreases (longer chains): · macro-adhesion increases (due to greater adsorption and more dissipative motions of chains) · macro-friction decreases (due to chain orientation phenomena) · nano-adhesion increases (due to chain adsorption on the AFM tip) · nano-friction increases (due to chain adsorption on the AFM tip). When free chains are present: · macro-adhesion decreases (due to a weak boundary layer effect, hindering the stress transfer) · macro-friction globally decreases · nano-adhesion increases (adsorption on the AFM tip, due to the great mobility) · nano-friction increases (adsorption on the AFM tip).

5.4 Conclusion

The experimental results emphasize the major role of molecular parameters such as free and pendant chains. Long and numerous free and pendant chains favor chain adsorption on the substrate. They also allow a greater dissipation during separation (tack experiments). However, these long chains can be oriented and extended during macroscopic friction, avoiding the direct transmission of stress to the network and decreasing the friction resistance. So the presence of free chains decreases the macroscopic adherence (they can form a weak boundary layer), and their effect on macroscopic friction is negligible. The nanoscale behavior is different: in that case the presence of free chains allows a better nano-adherence. These chains constitute a “viscous layer” that

References

can even trap the AFM tip. Nanoscale friction is also different, with a greater friction force for longer chains. Hence, nanoscale behavior is more sensitive to the adsorption ability due to molecular mobility. The effects of elastic contact and bulk dissipation are minimized and the interfacial adhesive behavior is magnified.

References 1 A. N. Gent, J. Schultz, J. Adhesion 1972, 3, 2 3 4 5

281–294. D. Maugis, M. Barquins, J. Phys. D: Appl. Phys. 1978, 11, 1989–2023. P. G. De Gennes, Langmuir 1996, 12, 4497–4500. H. Chun, A. N. Gent, J. Polym. Sci.: Part B: Polym. Phys. 1996, 34, 2223–2229. S. Bistac, J. Colloid Interf. Sci. 1999, 219, 210–211.

6 A. Galliano, S. Bistac, J. Schultz, J. Coll.

Interf. Sci. 2003, 265, 372–379. 7 A. Galliano, S. Bistac, J. Schultz, J. Adhe-

sion 2003, 79, 973–991. 8 T. Elzein, A. Galliano, S. Bistac, J. Polym.

Sci., Part B: Polym. Phys. 2004, 42, 2348– 2353. 9 S. Bistac, A. Galliano, Tribology Lett. 2005, 18, 21–25.

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6 Chemical Structure Formation and Morphology in Ultrathin Polyurethane Films on Metals C. Wehlack and W. Possart

Abstract

The polyurethane-metal interphase is studied in thin films of thicknesses varied within dPU = 15 nm – 2 lm on Au, Al and Cu and in the bulk with respect to chemical structure formation (FTIR) and morphology (OM, SFM). The results confirm that bulk properties generally dominate the integral behavior of some lm thick PU films on any metal, but even ultra-thin films on Au (some 10 nm) are in very good agreement with the bulk. In ultra-thin layers on Al and Cu, however, PU-metal interphase properties are clearly pronounced: Formation of urea or urea-like species and initial stages of dewetting phenomena are observed and increase as Al and Cu substrates are exposed to humidity prior to coating. Very specific chemical reactions (e.g. decomposition of carbodiimide and uretdione groups) in films on Cu during post-curing give evidence for the dissolution and mobility of chemically active Cu2+-ions within the whole film. Furthermore, quantitative results of the polyurethane cure substantiate considerable polyurethane formation catalysis for films on Cu, but slightly increased isocyanate conversion rates are also measured in ultra-thin films on Al and Au as compared to the bulk.

6.1 Introduction

The performance and long-term stability of adhesive joints and coatings are strongly affected by the chemical and structural properties of the interphase region situated between the uninfluenced adhesive polymer bulk and the adherend [1–16]. Unfortunately, this interphase is often hidden to the experimentalist, e.g., in closed joints. However, the interphase can be elucidated to a good approximation by studying thin polymer films [16–20] which are often referred to as “open joints”. Film thickness is varied in this approach to obtain information on microstructure gradients (e.g., in Refs. [1, 3, 4, 19, 20]). Bulk properties are expected to Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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6 Chemical Structure Formation and Morphology in Ultrathin Polyurethane Films on Metals

Fig. 6.1 Strategy for the characterization of the interphase by thin films with various thicknesses as compared with the bulk.

dominate the integral properties of thick polymer films, but, with decreasing film thickness, polymer–substrate interactions are supposed to affect the results more and more, as illustrated in Fig. 6.1. Using this strategy, it is possible to investigate how the curing process, the resulting network structure and even aging effects differ for adhesive films as compared with the bulk [19–24]. Such differences can result from a chemical or catalytic influence of the substrate, or from preferential adsorption and immobilization of adhesive components on the substrate, or the substrate could even initiate phase separation, crystallization, or any other preferential molecular arrangement. The spatial constraint due to the polymer–substrate and polymer–atmosphere interfaces could alter the molecular mobility and diffusion of species significantly as well – not only for ultrathin films but also for closed joints. How all these alterations to the microscopic processes affect the emerging macromolecular structure in a curing polymer adhesive is widely unknown, but there is much empirical evidence that they result in modified technical properties. Therefore, selected spectroscopic and microscopic techniques are applied to the characterization of chemical structure formation and morphology in thin polyurethane (PU) layers on Au, Al, and Cu: Infrared spectroscopy offers convenient access to the integral chemical properties of thin-film and bulk-like polymer samples, while optical (OM) and scanning force microscopy (SFM) allow detailed insights into homogeneity and topology.

6.2 Materials and Methods 6.2.1 Sample Preparation

Polyurethane bulk and thin-film samples are prepared from a two-part PU adhesive formed by polyaddition of Bayer Desmodur®CD (89% diphenylmethane4,4'-diisocyanate 1 with 11% uretoneimine triisocyanate 2 as additive) with a polyol mixture of Bayer Desmophen®2060BD (linear poly(propylene ether) diol 3) and Bayer Desmophen®1380BT (poly(propylene ether) triol 4) in the stoichio-


metric ratio of isocyanate (NCO) to hydroxyl (OHtriol/OHdiol = 80 : 20) groups. The investigations focus on adhesive interactions and structure formation of the polyurethane components and reaction products themselves. Therefore, the resulting model adhesive differs from commercially available formulations by the lack of any further additives or filler particles. The experiments are performed in inert gas atmosphere (argon or dried air with dew point = –70 8C) to avoid parasitic reactions of isocyanate groups with water and to guarantee reproducible properties of the polymer material.

uretoneimine

The PU curing reaction starts instantaneously with mixing of the components. After 15 min of thorough stirring, the resulting prepolymer is closely homogeneous. Bulk samples with film thickness dPU = 1 mm are cast directly from this prepolymer while thin films are prepared by spin-coating from prepolymer solution in THF. Thereby, different film thicknesses within the range of dPU = 15 nm to 2 lm are achieved for different prepolymer concentrations. As nearly ideal substrates, metal physical vapor deposition (PVD) coatings (ca. 100 nm of Au, Al, or Cu) on silicon wafers are used, thus providing not only a very flat and smooth topography with roughness rRMS < 1.5 nm but also guaranteeing a chemically very well-defined metal surface with native (hydr)oxide. After PVD, metal substrates are stored in dried air (Al: one day, Cu: four days, Au: some weeks) unless otherwise specified.

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6.2.2 Experimental Characterization

This well-suited sample geometry gives experimental access to the PU–metal interphase by FTIR external reflection absorption spectroscopy (ERAS; angle of incidence: 708; p-polarized light) and by microscopic methods (OM, tapping mode SFM). The investigations are performed in situ during curing (tcure = 0–72 h) at room temperature and on the fully cured state after 4 h of post-curing at 90 8C. ERA spectra of thin films on metals represent the integral chemical structure of the polymers as an average over the film thickness. The polymer properties in thin films are then compared with the bulk as measured by FTIR attenuated total reflection spectroscopy (ATR). A modest refractive index of the internal reflection element (IRE; made of ZnSe, n = 2.43 at 2000 cm–1) and an incident angle of 658 – still low but well above the critical angle of total reflection – are chosen for the p-polarized light in order to obtain an information depth of several microns. Hence, the ATR spectra provide the bulk properties of the polymer sample. Infrared spectroscopy experiments are performed on a Bruker IFS®66v/S FTIR spectrometer equipped with a Harrick Seagull® reflection unit. 6.2.3 IR Spectra Calculation

Inevitably, the evaluation of the thin-film ERA spectra has to account for the optical situation, which is significantly different from that for ATR bulk measurements. Spectra calculation provides the essential means for a reasonable comparison between bulk and thin-film spectra, thus allowing a detailed quantitative analysis [19]. In external reflection spectroscopy of thin films, the incident light undergoes partial reflection and refraction at the interface to air. These multiple reflections and refractions produce an interference pattern that modulates the spectrum in a characteristic way as a function of film thickness [25–27]. Furthermore, ATR spectra strongly depend on penetration depth as a function of wavenumber and refractive index of the IRE. As a result, neither ATR spectra with ERA spectra nor ERA spectra from polymer films of different thicknesses can be compared directly [19]. The problem of incomparability between bulk and thin-film spectra is solved in the following way (for the physics see, e.g., Refs. [25–27]). For the polyurethane bulk, the optical function given by Eq. (1) can be calculated from the ATR spectrum using software tools (e.g., Scout 2 [27]; based on matrix algorithm/Fresnel formulas and oscillator model parameter fit of the spectra). This was done with the ATR spectra measured for different polyurethane bulk states ^PU ~m for a given polyurethane both during and after curing. This experimental n state as well as the other optical parameters of a real thin-film sample experi^oxide ~m and thickness of the ment (angle of incidence, PU film thickness dPU, n ^metal ~m of the substrate metal) are used to calculate an underlying metal oxide, n


ERA spectrum for this sample. This calculated spectrum not only represents the bulk properties but also meets the necessary optical conditions of a given real ERAS experiment. Thus, it serves as a very good reference for measured ERA spectra of real thin-film specimens [19]. ^PU ~m; tcure nPU ~m; tcure i KPU ~m; tcure n

1

By using this method on any qualitative and quantitative FTIR results presented in this paper, thickness and substrate effects on the interphase are separated from the optical situation of the measurements. 6.2.4 IR Band Assignment

For a detailed interpretation of IR spectra, the following quantities of light absorption are required for every eigenvibration of the adhesive molecules: frequency, intensity, and – for a consideration of molecule orientation – the vector of the transition dipole moment. However, none of the available spectra catalogues and databases provides this complete information. They list only broad frequency ranges for IR bands and band intensities are just roughly classified. In most cases, rare chemical species are not considered. Furthermore, it remains uncertain how many eigenvibrations contribute to a given band in the measured IR spectrum and whether they are specific regarding relevant functional groups. These problems explain why the sole use of published databases does not provide a satisfactory and unequivocal band assignment in many cases. Therefore, we extend conventional IR band assignment by quantum mechanical (QM) modeling [28, 29]. In a first step, quantum mechanical modeling is used to obtain results for the molecular conformation and the geometry corresponding to the thermodynamically most stable state of the adhesive molecules [28]. Therefore, a force field optimization is applied on the manually set-up 3D structures using the molecular mechanics/dynamics calculation software Discover (Accelrys, San Diego). Several probable conformations are considered and the one with the lowest energy is chosen. Secondly, the density functional theory (DFT) calculation software DMol (Accelrys, San Diego) allows quantum mechanical refinement of the atom positions in the 3D structure. In a third step, DMol is used for the normal coordinate analysis that makes it possible to calculate all the normal vibration modes for these molecules, including the frequency and the vector of the transition dipole moment [28]. These results are then used for an extended and profound band assignment in the measured IR spectra. Thereby, the absorption index spectrum, as purely a property of the material, is preferred to reflectance or transmittance spectra, which are affected by the measuring conditions. As a result, procedures are developed that describe which IR bands can be used for peak analysis in order to monitor the curing reaction and the resulting network structure of the adhesive.

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Fig. 6.2 Selected normal modes of vibration for the uretoneimine group as examples for band assignment assisted by quantum mechanical modeling [29]. Light gray arrows: relative amplitudes of atom motion; dark

grey arrow: relative transition dipole moment. Atom code: H = white; C = light gray; O = dark gray; N = black. (a) ~meigen = 1859.0 cm–1; (b) ~meigen = 1722.5 cm–1.

In addition to the polyol and isocyanate components of the adhesive, relevant reaction product fragments/intermediates (e.g., urethane) and some side reaction products (e.g., carbodiimide, urea, uretdione, isocyanurate, etc.) are also considered. Selected examples of normal modes with respect to the uretoneimine additive are shown in Fig. 6.2. Band assignment and spectra interpretation in this paper are based on further, more detailed results given in Ref. [29].


First, attention will concentrate on the polyurethane curing reaction itself: that is, the polyaddition of isocyanate and hydroxyl groups. Degree of cure and isocyanate conversion rates (as deduced from FTIR) will be discussed for the bulk and compared with thin films on Au, Al, and Cu (Section 6.3.1). Secondly, the morphology of cured films (OM, SFM) will be treated in Section 6.3.2. Then, discussion will focus in more detail on the chemical structure (FTIR) of thin films on the different metals as compared with the bulk. In this way, substrate-specific effects on the interphase can be depicted for RT-cured and post-cured samples (Section 6.3.3). 6.3.1 Curing at Room Temperature

Fig. 6.3 depicts the evolution of IR band intensities during RT curing for the urethane group (e.g., 3311 cm–1, 1725 cm–1, or 1228 cm–1) that develops by con-


uretoneimine

uretdione

Fig. 6.3 IR spectra for the curing reaction of dPU = 1 lm polyurethane film on Au at room temperature.

sumption of isocyanate (2275 cm–1) and hydroxyl (3600–3300 cm–1) groups while methyl (e.g., 1457 cm–1 or 1376 cm–1), methylene, or ether species (1113 cm–1) remain constant according to Eq. (A).

R1 N C OH O R2

! R1

H O j jj N C O R2

A

In order to monitor the crosslinking, a quantitative analysis is performed for the intensity I2275(tcure) (= peak height) of the isocyanate band at 2275 cm–1 (asymmetric stretching) divided by the intensity I1376(tcure) (= peak height) of the methyl band at 1376 cm–1. This normalized band intensity I2275,norm(tcure) provides an IR spectroscopic degree of cure with respect to isocyanate groups according to Eqs. (2). UPU tcure

1

I2275;norm tcure I2275 tcure ; I2275;norm tcure I2275;norm tcure 0 I1376 tcure

2

77

78


In order to obtain I2275,norm(tcure = 0) as the initial state of the adhesive prior to curing, I2275,norm(tcure) is extrapolated. Fig. 6.4 illustrates the resulting degree of isocyanate conversion for the RT curing of 180 nm and 60 nm films on Al, Cu, and Au. For comparison, the conversion curve for calculated spectra is shown for the corresponding thin-film sample with bulk properties as it would have been measured under the given experimental conditions. The results substantiate increased isocyanate conversion rates in the vicinity of any of the three metals, even in ultrathin films on Au, to a greater extent on Al, and to the most pronounced extent on Cu (Fig. 6.4 b). As the film thickness

Fig. 6.4 Degree of isocyanate conversion (asym. stretching, 2275 cm–1) in ultrathin films on Al–, Cu– and Au–PVD during adhesive cure at room temperature. (a) dPU = 180 nm; (b) dPU = 60 nm.


is increased from 60 nm to 180 nm, bulk-like behavior is observed on Au whereas reaction acceleration on Al now becomes more moderate (Fig. 6.4 a). Obviously, considerable catalysis of polyurethane formation, raising the reaction rate by some orders of magnitude, is only found on Cu for films in this thickness range. This catalysis also becomes more emphasized with decreasing film thickness, i.e., with increasing dominance of interphase properties due to adhesive–substrate interactions. Except for ultrathin PU films on Cu, isocyanate conversion (2275 cm–1, e.g., in Fig. 6.3) is advanced but still incomplete after 72 h of cure at room temperature. Hence, post-curing is imposed on all samples in order to obtain fully cured polyurethane. 6.3.2 Morphology of Thin Films

Besides some minor thickness nonuniformity in thick (dPU > 400 nm) PU layers, the resulting spin-coating films are very smooth (RMS surface roughness is in the region of 1 nm; cf. Fig. 6.5 b: rRMS = 0.466 nm, z range = 6.485 nm) if metal substrates have been stored in dried air. As illustrated in Fig. 6.5 a by difference interference contrast, only a few singularities in topography can be found as a consequence of polymer shrinkage of about 5.4% in the bulk during RT curing. After a few hours of Al substrate storage in ambient air, however, ultrathin PU films tend to de-wet during curing as film thickness is decreased to dPU £ 100 nm (Fig. 6.6). Ultrathin film samples on Cu after several days of substrate storage in ambient air are subject to de-wetting, too. Films on Au, however, stay

Fig. 6.5 Morphology of a post-cured polyurethane film (dPU = 80 nm) on Al–PVD: Al stored for 1 d in dried air prior to coating ? smooth PU surface with few singularities. (a) Optical microscope image (DIC); (b) SFM height image (tapping mode).

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Fig. 6.6 Morphology of a post-cured polyurethane film (dPU = 80 nm nominally) on Al–PVD: Al stored for 1 d in humid ambient air prior to coating ? de-wetting of PU. (a) Optical microscope image (DIC); (b) SFM height image (tapping mode).

unaffected. Hence, the de-wetting should be induced by the changing oxide/hydroxide states on the surface of Cu and Al. Formation of typical de-wetting phenomena arises within the first few hours after spin-coating from the still liquid but, with increasing curing, more and more viscous PU prepolymer. These phenomena – ranging from waviness and a craterlike shape of the film to single droplets (e.g., as in Fig. 6.6 a, b) on the substrates – can be seen in the topography of RT-cured and post-cured films. They have been frozen by the competing network formation in their more or less early stages. 6.3.3 Chemical Structure of Cured Films

As intended by the adhesive formulation, no parasitic side reactions such as formation of urea, new uretdione, isocyanurate, biuret, or allophanate are detected for the bulk and for thick films on Al, Cu, and Au (dPU ³ 1 lm; cf. Figs. 6.3, 6.7). Even ultrathin PU films on Au correspond to bulk properties to a very good approximation (Fig. 6.8), which is a recommendation for films on Au as references with respect to effects in the interphase on different metals. However, spectra from ultrathin films on Al and Cu (Fig. 6.9, dPU = 55 nm) reveal very specific features. The curved baseline for Al confirms the initial stage of de-wetting as observed by microscopy; it results in a characteristic light-scattering topology (Fig. 6.9). New chemical species with an IR band at 1640 cm–1 are present in ultrathin films on Al after RT-curing. In ultrathin films on Al and Cu, increased phenylene band intensity at 1513 cm–1 on the one hand, and reduced phenylene intensity at 1615 cm–1 on the other, indicate either orientation phenomena or a chemical reaction affecting these eigenvibrations of the phenylene groups.


uretoneimine

uretdione

uretoneimine

uretdione

Fig. 6.7 Thick, post-cured PU films (dPU = 2 lm) on Al, Cu, and Au, as compared with the bulk.

Fig. 6.8 Bulk-like behavior of ultrathin, post-cured PU film (dPU = 80 nm) on Au.

81


uretdione

Fig. 6.9 Thin, post-cured PU films (dPU = 55 nm) on Al, Cu, and Au, as compared with the bulk.

uretoneimine

82

The extent of these specific effects grows with higher storage time of the native Al surfaces in ambient air and with decreasing film thickness. While only a few hours are sufficient for Al, it takes a few days of Cu substrate storage time in ambient air prior to PU coating before ultrathin PU films also tend to dewetting, to formation of a new peak at 1675 cm–1, and to phenylene intensity changes (cf. Fig. 6.10) as the film thickness is decreased. With respect to thinfilm and chemical structure formation, Al is obviously much more sensitive to humidity exposition than Cu. The extent of 1640 cm–1 or 1675 cm–1 band formation on the one hand, and phenylene peak intensity changes on the other, seem to be closely related, as one effect is never observed without the other. Further, there are connections with the availability of moisture during substrate storage, which increases the amount of adsorbed water and hydroxide on the surface that would react with the isocyanate groups. Therefore, urea or urea-like species (showing bands at 1695– 1630 cm–1 [30]) are the most reasonable candidates for these chemical structure modifications observed in the interphase on Al and – less intensely – on copper.


Fig. 6.10 RT-cured PU films (dPU = 80 nm) on Cu with different extents of exposure to humidity prior to coating.

uretoneimine m-NC,CO

Post-curing of the PU bulk and thick films on any metal leads to complete NCO conversion after 4 h at 90 8C in argon (Figs. 6.7 and 6.11). Films on Cu were already fully cured after 72 h at room temperature. Isocyanate bands also disappeared completely in thin post-cured films on Au whereas traces of residual isocyanate are still detected in ultrathin films on Al (Fig. 6.9). Some inhibition, reduced mobility, or steric hindrance obviously prevents isocyanate groups on Al from addition to still-available hydroxyl groups (3600–3400 cm–1). Moreover, urea peak and intensity changes of phenylene bands of the RT-cured PU on Al (and also on Cu) turn out to be irreversible with respect to post-curing. Further investigations are needed to reveal more accurately the chemical mechanisms underlying these phenomena.

Fig. 6.11 Consumption of residual NCO and formation of carbodiimide from uretoneimine during post-curing of the PU bulk.

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During post-curing, carbodiimide groups (doublet at 2100–2150 cm–1 in Figs. 6.7, 6.9, and 6.11) are produced on Al and Au and in the bulk by consumption of uretoneimine groups (characteristic of Desmodur®CD, 1700– 1730 cm–1 in Fig. 6.7, overlapping with urethane C=O stretching; ca. 1880 cm–1 in Fig. 6.11) according Scheme 6.1. O jj C R1

%

e

N e

R1

N C jj N

! R1

NCN

R1 O C N

%

R1

R

R2

1

O

H ! R2

O

O H jj j C N

R1

This reaction is associated with the formation of new isocyanate groups, which immediately add to excess hydroxyl groups to form more urethane species. The emergence of carbodiimide is still incomplete after 4 h of post-curing in films on Au and Al and in the bulk. Only on copper are no carbodiimide bands observed after post-curing, even in thick films (e.g., dPU = 2 lm in Fig. 6.7). Instead, a new chemical species develops with an IR band at 1660 cm–1, which is probably a reaction product of the carbodiimide consumption. Uretdione groups (isocyanate dimers; traces originate from the “as-received” Desmodur®CD component) are subject to similar effects. In films on copper, uretdion bands (1780 cm–1) vanish during post-curing, independently of film thickness (Figs. 6.7 and 6.9). The specific mechanism of carbodiimide and uretdione conversion on Cu is not sufficiently understood, yet. We still assume very specific reactions due to the influence of mobile copper ions. These findings can be related to the detection of Cu2+ ions by XPS and TOF-SIMS, even within thick polyurethane layers on native copper surfaces [20]. It should be noted that carbodiimide groups add water under humid conditions (e.g., in ambient air). Thus, urea groups are formed as a product of this reaction according to Eq. (B).

R

1

N C N R H O H ! R 1

1

H O H j jj j N C N R1

B

Therefore, carbodiimide can species no longer be detected after a couple of days in ambient air, as indicated in Fig. 6.12. Instead, a weak urea band is found at ca. 1650 cm–1. More details are given in Ref. [23]. The specific reactions of the uretoneimine additive and its derivatives are certainly not representative for polyurethane polymers in general. They still give further evidence for characteristic processes in the interphase on different metals. In addition, the mobility of the polymer network of this specific adhesive formulation and its mechanical properties will of course be affected by the associated substitution of each uretoneimine triple interlink by two linear connections.

6.4 Conclusion

Fig. 6.12 Formation of urea from carbodiimide in the polyurethane bulk in humid air (90% r.h., 40 8C).

6.4 Conclusion

The chemical structure and morphology of thin polyurethane films with various thicknesses have been characterized in order to gain insight into different formation and, different resulting microstructure of interphases on Au, Al, and Cu substrates as compared with the bulk. The results confirm that bulk properties generally dominate the integral behavior of thick (i.e., several microns deep) PU films on any metal. Additionally, even ultrathin films cured on Au have behavior in very good agreement with the bulk. The latter result proves that contributions of the PU–atmosphere interphase to the integral properties are negligible even for the very thin films. On Al and Cu, however, the PU–metal interphase properties are clearly pronounced in ultrathin layers (below 200 nm). Thus, the scientific approach to measure in-

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terphase properties by studying thin films proves to be particularly beneficial. New chemical species form, such as urea-like species on Al and Cu. Only on Cu are carbodiimide and uretdione groups consumed during post-curing; the mechanism is still under investigation at this stage of work. But even if the details are not yet fully understood, this mechanism clearly reveals the long-range impact of Cu substrates on the interphase formation, which is based on the dissolution and mobility of chemically active Cu2+ ions in the film. These copper ions may give rise to thermal oxidation processes or aging effects within the polymer. In general, the quantitative results of the polyurethane cure substantiate increased isocyanate conversion rates in the vicinity of any of the three metals, being most pronounced on Cu and to some extent on Al but also detected even in ultrathin films on Au. Nevertheless, considerable catalysis, by some orders of magnitude, of polyurethane formation is only found for films on Cu. Once more, the effectiveness of copper even in films of some microns thickness is easily explained by the dissolution of Cu2+ ions. Chemical structure and thin-film morphology on Al and Cu are very sensitive to substrate storage. Thin PU films are very smooth and homogeneous if prepared on substrates that were stored in dried air prior to coating. But if substrate storage conditions are humid, films on Al (after several hours) or Cu (after several days in ambient air) are more and more subject to de-wetting effects as their thickness is decreased. The resulting structures arise during curing within the first few hours after spin-coating and can be seen in the topography of cured films – “frozen” in their more or less early stages by the competing network formation. Some substrate-specific effects, as indicated above (such as the formation of urea-like species), are also more pronounced in films on Al or Cu stored in humid conditions. It can be concluded after all that the results reveal very specific features caused by metal–polymer interactions in the interphase of cured PU on different metal surfaces, though further experimental efforts are needed to gain a more complete understanding of the underlying chemical reactions and adhesion mechanisms. Furthermore, remarkable quantitative differences can be seen in the chemical structure formation of thin PU films with regard to reaction rate and degree of isocyanate consumption.

Acknowledgments

The financial support by the Bundesministerium für Bildung und Forschung (BMBF) is gratefully acknowledged. Further thanks go to Ms. Awa Sow (film preparation, FTIR) and Mr. Roland Krämer (SFM) for experimental assistance.

References

References 1 J. Bouchet, A. A. Roche, E. Jacquelin, J. 2 3 4 5

6

7 8

9 10

11 12

13 14

15

16

Adh. Sci. Technol. 2001, 15, 321–343. S. Bentadjine, R. Petiaud, A. A. Roche, V. Massardier, Polymer 2001, 15, 6271–6282. J. Bouchet, A. A. Roche, J. Adh. 2002, 78, 799–830. A. A. Roche, J. Bouchet, S. Bentadjine, Int. J. Adh. Adhes. 2002, 22, 431–441. J. Marsh, L. Minel, M.-G. Barthes-Labrousse, D. Gorse, Appl. Surf. Sci. 1998, 133, 270–286. S. Dieckhoff, R. Wilken, in Adhäsion und Vernetzung eines Modell-Epoxids auf oxidierten Aluminium- und Magnesiumoberflächen – I (Eds.: W. Possart, J. K. Krüger, S. Dieckhoff, T. Britz, H. Neurohr, B. Valeske, C. Wehlack, R. Wilken), DFG-project report no. Po 577/3–1, Kr 653/9–1, Df 792/1–1, July 2002, pp. 54–58. L. H. Sharpe, J. Adh. 1972, 4, 51 ff. C. Bischof, W. Possart, Adhäsion – Theoretische und Experimentelle Grundlagen, Akademie-Verlag, Berlin, 1983. M. Lotfipour, L. A. Reeves, D. Kiroski, D. E. Packham, J. Adh. 1994, 47, 33–40. R. Kraus, W. Wilke, A. Zhuk, I. Luzinov, S. Minko, V. Voronov, J. Mater. Sci. 1997, 32, 4397–4403. W. Possart, V. Schlett, J. Adh. 1995, 48, 25–46. J. Bouchet, A. A. Roche, E. Jacquelin, G. W. Scherer in Adhesion Aspects of Thin Films (Ed.: K. L. Mittal), Vol. 1, Brill Academic Publishers, Leiden, 2001. E. Mäder, K. Mai, E. Pisanova, Composite Interfaces 2000, 7, 133–147. W. Possart, J. K. Krüger in Proc. 25th Ann. Meet. Adh. Soc., WCARP II (Eds.: A. V. Pocius, J. G. Dillard), Orlando, FL, February 10–14, 2002, p. 104 ff. W. Possart, J. K. Krüger, Proc. EURADH 2002: 6th European Adhesion Conference and ADHESION ’02: 8th Int. Conf. on the Science & Technology of Adhesion and Adhesives, The Institute of Materials UK, IOM Communications, Glasgow, September 10–13, 2002, p. 7–8. W. Possart, S. Dieckhoff, D. Fanter, T. Gesang, A. Hartwig, R. Höper, V. Schlett, O.-D. Hennemann, J. Adh. 1996, 57, 227–244.

17 W. Possart, V. Schlett, J. Adh. 1995, 48,

25–46. 18 M.-G. Barthes-Labrousse, J. Adh. 1997,

57, 65–75. 19 C. Wehlack, W. Possart, Macromol. Symp.

2004, 205, 251–261. 20 S. Dieckhoff, R. Wilken, M. Noeske, Ad-

21

22

23

24

25 26 27

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häsions- und Alterungsmechanismen in Polymer–Metall-Übergängen (Ed.: S. Dieckhoff), BMBF-project report no. 03D0074, TIB Hannover, 2004. C. Wehlack, W. Possart, Chemical structure formation and morphology in ultrathin polyurethane films on metals, Proc. oral pres., 7th European Adhesion Conference EURADH 2004, Freiburg im Breisgau, September 5–9, 2004. C. Wehlack, A. Meiser, W. Possart, Chemical processes in ultrathin epoxy films on metals during ageing, Proc. oral pres., 7th European Adhesion Conference EURADH 2004, Freiburg im Breisgau, September 5–9, 2004,. C. Wehlack, R. Krämer, W. Possart, Ageing of a polyurethane adhesive: bulk versus ultrathin films on metals, Proc. poster pres., 7th European Adhesion Conference EURADH 2004, Freiburg im Breisgau, September 5–9, 2004. J. Kanzow, F. Faupel, W. Egger, P. Sperr, G. Kögel, C. Wehlack, W. Possart, Depthresolved analysis of the ageing behaviour of epoxy resin thin films by positron spectroscopy, Proc. oral pres., 7th European Adhesion Conference EURADH 2004, Freiburg im Breisgau, September 5–9, 2004. M. Milosevic, S. L. Berets, Appl. Spectrosc. 1993, 47, 566–574. O. Stenzel, Das Dünnschichtspektrum, Akademie-Verlag, Berlin, 1996. W. Theiss, Scout 2 – Spectrum interpretation by simulation, M. Theiss – Hardand Software for Optical Spectroscopy, Aachen, 2001. B. Schneider, Untersuchung von Adhaesiv/ Substrat-Grenzbereichen mit InfrarotSpektroskopie und Molecular Modelling, Dissertation thesis, Saarland University, Saarbrücken, July 2001. C. Wehlack, B. Schneider, M. Ott, W. Possart, IR spectroscopy of the polyurethane

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6 Chemical Structure Formation and Morphology in Ultrathin Polyurethane Films on Metals cure: band assignment assisted by quantum mechanical modeling, Proc. poster pres., 7th European Adhesion Conference EURADH 2004, Freiburg im Breisgau, September 5–9, 2004.

30 D. Lin-Vien, N. B. Colthup, W. G. Fateley,

J. G. Grasselli, The Handbook of Infrared and Raman Characteristic Frequencies of Organic Molecules, Academic Press, Boston, MA, 1991.

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7 Properties of the Interphase Epoxy–Amine/Metal: Influences from the Nature of the Amine and the Metal M. Aufray and A. A. Roche

Abstract

Epoxy–amine systems are used as adhesives and paints in many industrial applications. When they were applied on metallic substrates and cured, an interphase was created between the substrate and the polymer. The interphase had chemical, physical and mechanical properties, different than the bulk phase. Amines are known to chemically react with metallic ions to form organometallic complexes. Two amines were used: the diethylenetriamine (DETA) and the isophorone diamine (IPDA). When the IPDA organometallic complexes were over the solubility limit, chelates crystallized, whereas for DETA, organometallic complexes were formed but never crystallized. The crystals were analysed. For Al-IPDA crystals, the melting point was in the range of 75–80 8C. The new mechanical properties came from the crystals, that remained enclosed within the vitrified polymer; this was due to the crystals melting point, which was higher than the DGEBA-IPDA systems vitrification temperature.

7.1 Introduction

Epoxy–amine liquid prepolymers are extensively applied to metallic substrates and cured to obtain painted materials or adhesively bonded structures. Overall performances of such systems depend on the interphase created between the organic layer and the substrates. When epoxy–amine liquid mixtures are applied to a more or hydrated metallic oxide layer (such as Al, Ti, Sn, Zn, Fe, Cr, Cu, Ag, Ni, Mg, or E-glass), amine chemical sorption concomitant with metallic surface dissolution appear, leading to the organometallic complex or chelate formation [1, 2]. Furthermore, when the solubility product is exceeded, organometallic complexes may crystallize. These crystals induce changes of mechanical properties (effective Young’s modulus, residual stresses, practical adhesion, durability, etc.). Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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7 Properties of the Interphase Epoxy–Amine/Metal

In this work, two different amines were used and mixed with DGEBA (diglycidyl ether of bisphenol A) epoxy monomer (a/e = aminohydrogen function/ epoxy function = stoichiometric ratio = 1) to form bulk materials or coatings. The IPDA (isophoronediamine) is a cycloaliphatic diamine which may crystallize after modification (i.e., after being applied to the metal surface, leading to chemical reaction between liquid amine and metal), whereas the DETA (diethylenetriamine) is an aliphatic triamine which never crystallizes, even after 3 h in contact with aluminum. We evaluated some properties of the IPDA crystals (i.e., IPDA reacted with the metal, and crystallized), such as their melting point, and correlated the results with the change in mechanical properties. Finally, we studied the influence of the stoichiometric ratio a/e on the polymerization of our systems.

7.2 Materials and Methods 7.2.1 Materials

The metallic substrates used were 1.200 ± 0.005 mm thick 1050 (0.25% Si, 0.4% Fe, 0.05% Cu, 0.05% Mn, 0.05% Mg, 0.05% Zn, 0.03% Ti, 0.05% V, i.e., min. 99.5% Al) commercial aluminum alloys from Péchiney. Aluminum sheets were prepared by die-cutting to provide identically sized strips (50 ´ 10 mm2). Before any polymer application, aluminum substrate surfaces were cleaned by ultrasonic immersion in acetone for 10 min, wiped dry, submerged in a sulfochromic solution (250 g L–1 of sulfuric acid (d = 1.84), 50 g · L–1 of chromium (VI) oxide and 87.5 g · L–1 of aluminum sulfate octadecahydrate) 1 h at 60 8C, rinsed in running water for 1 min, allowed to stand in deionized water for 5 min and wiped dry. After surface treatment, all substrates were stored less than 2 hours in an airconditioned room (20 ± 2 8C and 50 ± 5% r.h.), before polymer application. The epoxy prepolymer used was diglycidyl ether of bisphenol A (MW = 348 g mol–1, DGEBA DER 332 from Dow Chemical). The curing agents were either IPDA from Fluka or DETA from Aldrich. Assuming a functionality of 4 for IPDA, 5 for DETA, and 2 for the epoxy monomer, the stoichiometric ratio a/e used was equal to 1 (exceptions are mentioned). To control the extent of chemical reactions between the metallic surface and liquid monomers, leading to the formation of a thick interphase, liquid epoxy– amine mixtures were kept in contact with the metallic surface at room temperature for various periods of time before the desired adhesive curing cycle was started (e.g., see Fig. 7.1): at 190 8C, vitrification appeared within a few minutes, stopping any reaction between amine and metal, and/or diffusion phenomenon. These curing cycles allowed the maximum conversion (i.e., the maximum glass transition temperature). Conversely, when interphase formation was not desired,


Fig. 7.1 Curing cycles of the DGEBA–IPDA and DGEBA–DETA systems.

coated specimens were placed in the preheated oven immediately after the epoxy–amine application (i.e., within less than 1 min). 7.2.2 Thermal Analysis (DSC)

Differential scanning calorimetry (DSC) experiments were carried out in a Mettler (DSC 30) apparatus to determine the onset glass transition temperature (Tg onset) of epoxy resins. Sealed aluminum pans containing 15–20 mg of resin were heated from –50 8C to 250 8C at a rate of 10 K min–1 under a continuous flow of U-grade argon. Samples were weighed using a Mettler balance having a ±5 lg sensitivity. The calorimeter was calibrated with both indium and zinc. The glass transition temperature was determined with ±1 K accuracy. 7.2.3 Micro-Infrared spectroscopy (l-FTIR)

“Micro-IR” maps were made using FTIR Imaging Spotlight 300 from Perkin ElmerTM. To determine the practical adhesion, thick stiffeners (25 ´ 5 ´ 4 mm3) made of polymer were molded onto the metallic substrate and debonded by mechanical testing with a three-point flexure test (ISO 14679). After this test, it was possible to cut 1 mm thick slices of polymer (perpendicularly to the adherend surface) and to analyze them. A transmission infrared map could be realized, data points being collected at every 6 lm interval of sample displacement along a line perpendicular to the metal surface. According to the D66545 Perkin-ElmerTM Product Note, the dual imaging resolution was 6.25 lm pixel size. Infrared spectra were recorded in the 3000–7800 cm–1 range using a dual-mode

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detector. The imaging mode was used. For each analysis, 128 scans were collected at 4 cm–1 resolution. 7.2.4 Fourier Transform Near-Infrared Spectroscopy (FT-NIR)

A near-infrared spectrometer (Equinox 55 from Bruker) was used with OPUS software. The near-infrared spectra were recorded in the 4000–8000 cm–1 range. For transmission analysis, polymer coatings (from 50 lm to 2 mm thick) were analyzed after debonding of the coating from the substrate by bending the coated substrate, and background spectra in air were recorded. For each analysis, 128 scans were collected at 4 cm–1 resolution. 7.2.5 Inductively Coupled Plasma Spectroscopy (ICP)

An ICP spectrometer (simultaneous Vista from Varian) was used with a 40 MHz generator and a new CCD detector (70 908 pixels, wavelength from 167 to 785 nm) to determine the metal ion concentration within the liquid amine after a 3 h liquid/solid contact time. Distilled water was used as the diluting agent (each sample – from 1 to 10 mg – was diluted in 50 mL of distilled water). 7.2.6 X-Ray Diffraction (XRD)

The structural studies of the organometallic crystals had been conducted at room temperature by X-ray diffraction (XRD). Considering their initial shape (sharp needles) and dimensions (1 lm in diameter and 50 lm long), it was necessary to increase their size, so they were dissolved in deionized water and after a very slow water evaporation process, the organometallic complexes recrystallized. The shape factor was identical, but they were bigger (about 1 mm long). As they were single crystals (their extinction was observed using polarized optical microscopy, POM), they could be analyzed by XRD. Single-crystal diffraction data were collected using a Nonius Kappa CCD diffractometer. 7.2.7 Polarized Optical Microscopy (POM) Coupled with a Hot Stage Apparatus

Crystallized complexes or epoxy-modified amine mixtures could be melted or cured between two glass plates using a Mettler FP 82 hot stage, coupled with an FP 90 central processor, under POM (Laborlux 12POLS from Leica, equipped with a CCD IRIS color video camera from Sony). Crystals were heated from 25 to 100 8C at a rate of 1 K min–1 and mixtures were cured from 20 to 250 8C (see Section 7.3.4).



When liquid epoxy–amine prepolymers were applied and cured on metallic substrates, interphases were created within the organic layer in the vicinity of the metal surface. 7.3.1 Interphase Formation Mechanisms

In previous studies [3, 4], we had pointed out that the interphase formation mechanisms result from dissolution of the metallic surface layers, concomitantly with ion diffusion through the liquid prepolymer. In order to detect the dissolution phenomenon, pure amine (either DETA or IPDA) was previously applied to chemically etched metallic sheets (either Al or Ti alloys were used, and had hydroxidic surfaces). After 3 h, the metallic surfaces were scraped with a PTFE spatula. The “modified” amine (i.e., the amine reacted with the metal) was analyzed. Whatever the natures of the amine and the metal were, metal ions were detected in the “modified” amines by ICP analysis and new peaks were detected by infrared spectroscopy [5]. To indicate hydroxide dissolution, a very thin layer of liquid amine was applied to chemically etched aluminum, and Infrared Reflection – Absorption Spectroscopy (IRRAS) spectra were recorded every 5 min (the hydroxide band intensity variation at ca. 3430 cm–1 was followed). The OH group peak intensity decreased when the amine–metal contact time increased [5]. Conversely, if pure DGEBA monomer was applied to the metal surfaces, even after 3 h in contact with the metallic surfaces, no metal ion was detected by ICP in the DGEBA recovered, and the infrared spectra remained identical before and after the contact with the metal. Finally, if pure amine monomer was applied to gold-coated substrates, no chemical reaction was observed (by either ICP or FTIR analyses). The amine chemisorption onto oxidized or hydroxidized metal surfaces, concomitantly with the partial dissolution of the surface oxide (and/or hydroxide) metal substrate, was observed according to the fact that amine monomers are basic. Then it could be assumed that either: · metal ions diffuse within the liquid monomer mixture and react with amino groups of the hardener to form organometallic complexes, or · organometallic complexes form on the metal surface layer [6] and diffuse within the liquid monomer mixture (epoxy–amine). For all systems except those on gold, after application to metallic surfaces amines were modified to form an organometallic complex. Then, the initial liquid epoxy–amine mixture was transformed into a mixture of organometallic complexes and pure (i.e., unmodified) amine with pure epoxy prepolymer.

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7.3.2 Formation of New Networks

During the curing cycle, only uncrystallized organometallic complexes can react with the epoxy monomer, leading to a new network having a lower Tg. Previously, the presence of the two networks (the initial one and the “modified” one) was revealed by DSC and DMTA [3]. In order to detect modification of the network, some coatings were made on various kinds of substrate (Al, Ti, Sn, Zn, Fe, Cr, Cu, Ag, Ni, Mg, Au, and E-glass) covered by their natural oxide layer. Curing took place on the metals and on the E-glass. Then the coated substrates were bent to pull off the coating layer. The glass transition temperature could easily be determined for such free coatings. The relevant glass transition temperatures of coatings were compared with the bulk data (e.g., see Table 7.1). Whatever the amine (IPDA or DETA) and the metal were, the glass transition temperatures of the coatings were lower than that of the pure bulk Tg (except with a gold substrate). So, the mechanisms described previously were also observed for any metallic substrates provided they were covered by an oxide or hydroxide layer. 7.3.3 Crystallization of “Modified” IPDA

If the complex (or chelate) concentration within the liquid amine or epoxy– amine prepolymer was higher than its solubility limit, complexes (or chelates) crystallized. Sharp needle-like crystals were observed with modified IPDA what-

Table 7.1 Influence of the nature of the amine and the metal on the glass transition temperature Tg of coatings.

Metal

Bulk Cu Sn Cr Ag Si Fe Al 5182 Al 1050 Ni Ti Zn Mg Au

Tg [8C] DETA

IPDA

132 92 88 35 86 91 89 94 94 89 91 85 76 132

163 117 131 122 114 129 130 115 126 127 130 129 128 163


Fig. 7.2 Series of POM images showing determination of the melting point of crystals.

ever the nature of the metal (except gold), whereas DETA never crystallized (even after 3 h in contact with any metal). Considering their initial dimensions, it was necessary to increase their size. The organometallic complexes were therefore recrystallized. Then, crystals could be melted under POM (e.g., see Fig. 7.2) to determine their melting temperature and to observe, or not, extinction under polarized light. For Al–IPDA crystals, the melting point was in the range 75–80 8C, and the crystals observed were single crystals (extinction was observed using POM). They were analyzed by XRD and seemed to be orthorhombic (a = b = c = 908). 7.3.4 Modification of Mechanical Properties

We have mentioned that the crystal melting point was around 8 0 8C (see Section 7.3.3), but the highest temperature of the DGEBA–IPDA curing cycle was 190 8C (e.g., see Fig. 7.1), and we firstly mentioned that these crystals induced changes of mechanical properties, such as Young’s modulus, residual stresses, practical adhesion, and durability [7] (Section 7.1). In order to confirm the presence of crystals within the cured material, even after the entire curing cycle (so that they could act as short fibers after cooling), we followed their behavior during the DGEBA-“modified” IPDA curing cycle, using POM (e.g., see Fig. 7.3). Whatever the temperature was, needle-like shapes (like our initial organometallic crystals) could be observed. For our DGEBA–IPDA pure prepolymer mixtures, the vitrification temperature was 60 8C [3], so this DGEBA–IPDA system vitrified before the melting temperature (80 8C) of the crystals. For temperatures higher than 60 8C, even if organometallic complexes were melted (i.e., liquid), they remained trapped within the vitrified matrix. As the refractive indices of the vitrified polymer and the melting crystals were different, the interface formed between the liquid phase of the organometallic complex and the vitrified

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Fig. 7.3 Series of POM images showing evolution during a DGEBA/Al-modified IPDA curing cycle.


matrix could be observed during the full curing cycle. After the cooling down (Fig. 7.3, POM image 10), we were unable to tell whether the observed shapes showed any recrystallization of complexes or the presence of organometallic complexes remaining in their liquid state. However, according to their shape factor and their orientation (parallel to the metallic surface), only crystals could act as short fibers in the matrix and modify the polymer’s mechanical properties, as described in previous work [7]. For coatings 100 lm thick, longitudinal Young’s moduli were determined (using the three-point flexure test [8]) and found to be 5 and 2 GPa for DGEBA–IPDA and DGEBA–DETA coatings respectively. The Young’s moduli of the relevant bulk materials were about 3 GPa. As crystals were found only when IPDA hardener was used, the increase in Young’s modulus could be associated with the crystal formation and orientation [9]. 7.3.5 Comparison of Coatings and Metal–Bulk Interphases

To determine the practical adhesion, a polymer block 4–5 mm thick was molded onto the metallic substrate. The cured polymer stiffener was debonded during the mechanical test (ISO 14679). To verify that the interphase formation mechanisms took place either in the coatings or in the bulk, after the three point flexure test it was possible to slice the epoxy–amine polymer block (perpendicular to the surface of the adherend) and to analyze it. Then, a transmission infrared map could be realized. We considered the bands at 4530 cm–1 (epoxy combination), 6500 cm–1 (amine), and 4623 cm–1 (the aromatic C–H ring stretch combination, used as the reference) [2]. Normalized amine and epoxy band intensity variations are derived from l-FTIR spectroscopy for DGEBA–IPDA or DGEBA– DETA systems (e.g., see Fig. 7.4). The normalized amine band intensity variations are given for the IPDA system, whereas the normalized epoxy band intensity variations are shown for the DETA system. Both the normalized amine band intensity for the DETA system and the normalized epoxy band intensity for the IPDA system remained quite constant whatever the thickness, and were not represented. The top of the sample was initially in contact with the metal. On both maps, an interphase (corresponding to the region where band ratios vary), and a bulk region (with homogenous properties) were observed. The interphase thickness is about 300 lm for IPDA systems and 500 lm for DETA systems. In order to determine whether coatings and bulk materials behave identically near the metal surfaces, the l-FTIR analyses for polymer slices were compared with FT-NIR data from coatings of varying thickness. Amine (va) and epoxy (ve) degrees of conversion were calculated, using the ratio of the respective band areas, by Eq. (1).

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Fig. 7.4 Micro-FTIR map of the interphase for DGEBA–IPDA (left) and DGEBA–DETA systems (right).

veNIR 1

A4530 A4623 t A4530 A4623 t0

vaNIR 1

A6500 A4623 t A6500 A4623 t0

1

va and ve were determined for both systems (initially DGEBA–IPDA and DGEBA–DETA were mixed in the stoichiometric ratio (a/e = 1), then they were applied to the metal surface and state 3 had room temperature before the curing cycle was started) and photographed (e.g., see Fig. 7.5): va for DETA and ve for IPDA were quite constant and equal to 1, whatever the coating thickness. In addition, as observed by l-IR spectroscopy (e.g., see Fig. 7.4), va decreased near the metallic surface for the IPDA, while ve decreased near the metallic surface for the DETA system. The only difference between the two systems was the crystallization of organometallic complexes for the IPDA system, whereas DETA complexes never crystallized. Whatever the amine was (either IPDA or DETA),


Fig. 7.5 Degrees of conversion for amine (va) and epoxy groups (ve) versus coating thickness by NIR spectroscopy.

modified and unmodified amines had very similar NIR spectra [3]. In particular, we could find the same amine band at 6500 cm–1. Indeed, when modified IPDA crystallized, modified amines within the crystal phase could not react with epoxy groups, so unreacted amino groups were detected by NIR. In addition, epoxy groups might react together by an etherification reaction [3, 10] during curing, catalyzed by metallic ions. Consequently, all oxiranes could be converted. Conversely, modified DETA never crystallized. Assuming that the functionality of modified DETA was lower than pure DETA functionality (as explained in Section 7.3.6), all the amino groups that reacted with epoxy groups and epoxy groups in excess were detected by NIR. Epoxy groups could not react together by an etherification reaction which would need a higher temperature. DGEBA–

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DETA systems vitrify at room temperature in less than an hour, contrary to DGEBA–IPDA systems, which vitrify at 60 8C. When we cured the DGEBA– DETA systems (after 3 h at room temperature), they were already vitrified, preventing any etherification reaction. 7.3.6 Influence of the Stoichiometric Ratio

Finally, the variation of the glass transition temperature for both systems (DGEBA–IPDA or DGEBA–DETA) versus the stoichiometric ratio (a/e) is reported (e.g., see Fig. 7.6) for either pure or “modified” materials. Usually, as the functionalities of epoxy and amine monomer were well defined, mixing materials at the stoichiometric ratio of 1 led to the formation of the most crosslinked network having the highest glass transition temperature. From Fig. 7.6 it can be

Fig. 7.6 Variation of glass transition temperature as a function of stoichiometric ratio.

7.4 Conclusion

seen that functionalities of 4 for IPDA, 5 for DETA, and 2 for the DGEBA epoxy monomers are correct (the maximum glass transition temperatures were observed for a/e = 1, for both IPDA and DETA systems). But, for IPDA and DETA “modified” bulk materials, the maximum Tg values were observed for a stoichiometric ratio of 1.15. Indeed, after modification (i.e., formation of the organometallic complexes), the functionality of the organometallic complex was lower than the pure amine one. This assumption is in good agreement with findings on the l-FTIR maps (see Section 7.3.5). Using those systems (DGEBA and IPDA or DETA on metallic substrates), the stoichiometric ratio had to be higher than 1 in order to obtain the highest crosslink density (i.e., the maximum glass transition temperature) and the best material properties [9]. Moreover for a/e £ 1.15, the Tg values are always lower after modification than before (i.e., for pure materials). This decrease could be induced by a plasticization effect, due to small molecules such as unreacted organometallic complexes and/or water molecules diffusing within the liquid prepolymer following dissolution of the metallic hydroxide. In addition for a/e £ 1, the epoxy groups are in excess. They might react together by an etherification reaction [10] to form a less reticulated network, with a lower Tg. Using DGEBA–IPDA systems and varying the stoichiometric ratio from 0.7 to 1.3, Bentadjine [3] has shown that etherification decreases when the stoichiometric ratio increases and that the etherification phenomenon increases for DGEBA–IPDA coatings 100 lm thick on titanium alloy. In all our systems the corresponding ether bands (1120 cm–1) were also found.

7.4 Conclusion

When epoxy–amine prepolymers were applied on metallic substrates, interphases between the coating part, having the bulk properties, and the metallic surface were created. Amine chemisorption onto oxidic or hydroxidic metallic surfaces, concomitantly with partial dissolution of the surface oxide (and/or hydroxide) on the metal substrate, was observed, according to the basicity characteristics of the amine monomers (pKa ³ 10). Then it could be assumed that either: · metallic ions diffused within the liquid monomer mixture and reacted with amino groups of the hardener to form organometallic complexes, or · organometallic complexes formed on the metallic surface layer [6] diffused within the liquid monomer mixture (epoxy–amine). Whatever the amines were, after application to metallic surfaces they were “modified” to form organometallic complexes having a lower functionality than the pure amines. These complexes might crystallize if their solubility limit was exceeded. During the curing cycle, uncrystallized organometallic complexes reacted with the epoxy monomer leading to a new network having a lower Tg, whatever the amine and the metal were. For IPDA systems, however, needle-like

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structures were observed during all the curing cycles. They modify the polymer’s mechanical properties by acting as short fibers. Even if crystals melted at 80 8C, they remained embedded into the vitrified polymer. Finally, thin coatings and bulk materials formed (for ISO 14 679 mechanical tests) on metallic surfaces had the same interphase formed close to the metal.

Acknowledgments

We acknowledge the Perkin-ElmerTM Instruments team from Lyons (and particularly Dr. P. Delmont), for their help in making the l-FTIR maps. Also, we thank the crystallographic group of Claude Bernard University from Lyons 1 (principally D. Merle, Prof. Perrin, and Dr. Thozet) for doing XRD analyses and for their helpful discussion.

References 1 A. A. Roche, J. Bouchet, S. Bentadjine,

2

3 4

5

International Journal of Adhesion and Adhesives, 2002, 22, 431–441. S. Bentadjine, R. Petiaud, A. A. Roche, V. Massardier, Polymer, 2001, 42, 6271– 6282. S. Bentadjine, Ph.D. thesis, Villeurbanne, 2000, p. 157. M. Aufray, A. A. Roche, 27th Annual Meeting of the Adhesion Society, Wilmington, NC, 2004, pp. 367–369. A. A. Roche, M. Aufray, Swissbonding, 2003, 2–11.

6 F. Debontridder, Ph.D. thesis, Paris,

U.F.R. Scientifique d’Orsay, 2001, p. 223. 7 M. Aufray, A. A. Roche, 7th European Ad-

hesion Congress (Euradh), Freiburg, Germany, 2004, pp. 55–60. 8 M. Benabdi, A. A. Roche, J. Adhesion Sci. Technol., 1997, 11(3), 373–391. 9 J. Bouchet, A. A. Roche, P. Hamelin, Thin Solid Films, 1999, 355/356, 270– 276. 10 J. P. Pascault, H. Sautereau, J. Verdu, R. J. J. Williams, Thermosetting Polymers, Marcel Dekker, New York, 2002.

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8 Mapping Epoxy Interphases M. Munz, J. Chung, and G. Kalinka

Abstract

The mechanical characterization of epoxy interphases in composite systems is addressed. Beyond furthering the understanding of interface-related effects in thermosets, knowledge of interfacial stiffness variations is of major relevance for the calculation of mechanical stress distributions and the prediction of composite failure. In recent years, nano- and micro-indentation techniques which allow the investigation of cross-sectional surfaces have become available. In the case of narrow interphases, scanning force microscopy (SFM) can be employed for stiffness mapping, whereas depth-sensing micro-indentation (DSI) proves useful for wider interphases. Potential measurement artifacts due to changes of the tip–sample contact area and the lateral extent of the subsurface stress distribution beneath the SFM tip are discussed. Two different examples of epoxy interphase characterization are given. Epoxy interphases in copper/epoxy composites were investigated using force modulation microscopy (FMM, an SFMbased technique), and the epoxy interphase of a thermoplastic/epoxy composite was characterized by means of DSI. The thermoplastic was polyvinylpyrrolidone (PVP). In the copper/epoxy case, the epoxy stiffness was found to be reduced within the interphase, the total width of which was no more than about 29 lm. In the PVP/epoxy case, however, the stiffness gradient was extended over about 175 lm, and within the interphase the stiffness was found to be enhanced. The epoxy resin of the diglycidyl ether of bisphenol A (DGEBA) type was cured with the aromatic amine diaminodiphenylsulfone (DDS) at 170 8C. Across the interphase a depletion of DDS was detected by energy-dispersive analysis of X-rays (EDX). Within the substoichiometric regime of the epoxy system used, reduced amine concentrations were found to correspond to increased stiffness values.


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8 Mapping Epoxy Interphases

8.1 Introduction

By its very nature, the performance of adhesive joints is strongly affected by interfacial interactions. The mean conformation of the macromolecules located in the immediate vicinity of the interface of a polymer adhesive with the adherend is expected to be different from that of macromolecules located within the bulk of the adhesive. Physical or chemical adhesion interactions between segments of the macromolecules and the surface of the adherend will result in complex conformations such as trains of adhesion sites with narrow loops in between [1]. Similarly, adsorption of end groups can occur, leading to graft-like conformations. Such interfacial conformations result in variations of the local density as measured in the direction perpendicular to the surface of the adherend. Thus, adsorption and the corresponding changes of the mean conformation imply constraints to molecular motions which in turn find their expression in the relaxation behavior and local elastic properties [2]. The loss of entropy associated with interface-induced reduction of conformational freedom may also lead to a segregation in molecular weight. Assuming that no interface-induced ordering processes exist which generate a long-range order (such as transcrystallization), these effects are restricted to a region extending over a distance in the order of the radius of gyration of the macromolecules, i.e., some tens of nanometers at most [3]. Owing to the nonzero thickness of this interfacial zone with properties different from the bulk adhesive, it represents a volume rather than a mere area and usually it is referred to as an interphase (in the following denoted as IP). In the case of a thermosetting adhesive polymerized in the presence of the adherend, much longer-ranging effects can occur due to interface-induced diffusion [4–7]. For instance, when starting from a liquid mixture of a two-component system containing a prepolymer and a curing agent, preferential adsorption of one of these components onto the adherend surface will result in concentration gradients. A thermosetting system of considerable practical relevance is epoxy resin crosslinked with an amine curing agent. Owing to the polymerization and crosslinking reaction, spatial variations of the amine/epoxy concentration ratio will be converted into corresponding variations of the crosslink density which in turn affects the low-strain as well as the high-strain mechanical properties, such as elastic stiffness and strain-to-failure, respectively. Thus, given a sufficient driving force for interface-driven segregation, epoxy IPs are expected to extend over length scales given by diffusivity and the available time interval (see below) rather than solely the radius of gyration of the macromolecules. The diffusivity changes in the course of the curing reaction. With ongoing curing reactions, the average molecular weight increases, leading to a corresponding rise in the viscosity of the liquid mixture of epoxy resin and amine curing agent. Finally, at the point of gelation the network of crosslinks extends over macroscopic dimensions and the viscosity diverges. Hence it can be inferred that the time window for diffusion effects is a fraction of the gelation time, which in turn depends on the curing temperature and the peculiarities of the epoxy–

8.1 Introduction

amine system. Correspondingly, IP formation is controlled by the ratio between the characteristic times of diffusion and of curing reaction [5]. In principle, this ratio can also affect the IP concentration profile in a qualitative sense. When considering a homogeneous epoxy–amine mixture just brought into contact with the surface of an adherend or a filler particle, preferential adsorption of the amine molecules will result in local variations of the amine/epoxy concentration ratio r (to be defined below). Assuming a comparably fast diffusion of the amine molecules, their quick enrichment at the interface will result in a near-interface zone of increased r values and an adjacent zone of amine depletion, i.e., with reduced r values. Similar concentration variations are dealt with in the field of surface-driven phase separation in polymer solutions and mixtures [8]. While the wetting layer is in local equilibrium with the depletion layer, the diffusion from the bulk down the concentration gradient into the latter feeds the growth of the former. In the case of a high reaction rate of the epoxy system, gelation will be reached soon and the IP concentration profile encompassing a near-interface zone of amine enrichment and an adjacent zone of amine depletion will be “frozen” owing to the divergence of the viscosity. A corresponding composition perturbation near the air surface was reported on by Yim et al. [9], who performed neutron reflectivity measurements with thin epoxy films. The magnitude of the segregation effect was found to be a function of composition and cure temperature. In the opposite case of a low reaction rate, the amine depletion zone can be refilled from the bulk and the final IP concentration profile is likely to consist only of the amine enrichment zone. In this case a monotonic concentration profile is established with the amine concentration decaying from the interface value to the lower value of the bulk epoxy. The IP concentration gradients once established will translate into spatial variations of the final network of crosslinks as well as of the related mechanical properties. In a simplified approach, this kind of IP can be thought of as a series of slabs each of which is characterized by a different but constant concentration ratio r. One of the key parameters defining the mechanical properties of the cured epoxy is the crosslink density. In a diepoxy fully cured with a diamine, if additional effects such as side reactions are neglected, each primary amine group is expected to react with two epoxide groups. Branched crosslinks result from this situation. With the chemical functionalities fa = 4 and fe = 2 of the diamine and the diepoxy molecules, respectively, the amine/epoxy mixing ratio, r, is given by Eq. (1), where Na and Ne denote the respective numbers of moles [10]. r

N a fa ; N e fe

1

For a stoichiometric mixture of diamines and diepoxy, Na = 1 and Ne = 2. However, deviation from the stoichiometric situation results in either excess epoxide (r < 1) or excess amine groups (r > 1). Unreacted groups at terminal positions act

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as dangling ends. In analogy to thermoplastic polymers, an increasing average concentration of chain ends implies a higher free volume and a reduction of the epoxy Tg value [11]. In addition to the crosslink density, the epoxy stiffness depends on the packing density also. The stiffness of the adhesive and its IP variations are parameters of major relevance to the mechanical stresses occurring at interfaces. Having calculated IP stress distributions within arrays of reinforcing fibers embedded in a polymeric matrix, Anifantis reported on the dependence of step and peak gradients on the particular stiffness profile as measured in a radial direction [12]. Stress peaks can be considered as the likely positions of fracture initiation, once a certain stress level is exceeded. Hence, for the purpose of reliable assessment of interfacial stress distributions, a precise knowledge of IP property profiles is highly desirable. “The lack of knowledge of quantitative IP properties” was identified by Lesko et al. as “the Achilles’ heel of the field” [13]. However, in recent years nano- and micro-indentation techniques have become available which allow force penetration experiments to be performed with very sharp tips and at precisely controlled positions. Although there is a general acceptance of the relevance of IP properties, there still exist only a limited number of investigations providing quantitative information on mechanical IPs, including parameters such as IP width, total change in modulus across the IP, or radial stiffness profile [14–18]. In general, the available indentation techniques are based either on scanning force microscopy (SFM) or on dedicated indentation setups providing well-defined tip geometries. The two approaches are described in Sections 8.2.1 and 8.2.2, respectively. Issues related to potential artifacts in IP characterization are discussed in Section 8.3. Finally, examples of epoxy IP characterization by means of SFM-based stiffness mapping as well as depth-sensing micro-indentation (DSI) are given in Sections 8.4.1 and 8.4.2, respectively.

8.2 Stiffness Mapping by Indentation Techniques 8.2.1 SFM-Based Stiffness Mapping in Force Modulation Microscopy (FMM) Mode

A promising approach is the application of nano-indentation and related techniques, i.e., the measurement of force–penetration curves using very sharp tips and positioning elements with a precision on the nanometer scale. Scanning force microscopes are frequently used for such experiments, providing the ability to map spatial variations of properties such as adhesion or stiffness. In a straightforward approach, cross-sections of the composites to be analyzed can be prepared and the resulting surfaces studied using SFM-based stiffness mapping [15, 18]. The core element of a scanning force microscope is a micro-cantilever which carries a sharp tip (Fig. 8.1 a). When operating in contact mode, the

8.2 Stiffness Mapping by Indentation Techniques

Fig. 8.1 Schematic representation of the core elements of (a) a scanning force microscope, (b) the contact between tip and sample, and (c) the mechanical description in terms of a series of two springs. The two springs, kc and kts, represent the stiffness of

the cantilever and the tip–sample contact, respectively. The characteristic parameters of the tip–sample contact are the radius of curvature of the tip apex R, the contact radius a, and the deformation d.

apex of the tip is in permanent mechanical contact with the sample surface under investigation. A force in the direction perpendicular to the surface is exerted via the bending of the cantilever (to be exact, as a result of the slight cantilever tilt of *10–148 the force is not perfectly perpendicular to the surface). Typically, the radius of the apex of the tip, R, is in the range 10–50 nm. Depending on the local stiffness of the surface and the tip material, a deformation occurs at the tip–sample contact. The stiffness, kts, of the latter is proportional to the radius, a, of the contact area as well as to the reduced modulus, Ets (Fig. 8.1 b). According to the Hertz model, kts is given by Eq. (2) (see Refs. [19] and [20], and references therein). kts 2aEts

2

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The reduced modulus, Ets, is given in terms of the Young’s modulus values of both the tip, Et, and the sample material, Es, by Eq. (3), where mt and ms denote the Poisson’s ratios of tip and sample material, respectively [20]. 1 3 1 m2t 1 m2s ; Et Es Ets 4

3

A prerequisite for the application of the Hertz or Johnson–Kendall–Roberts (JKR) contact models is that the contact radius a is much smaller than the radius of curvature R of the apex of the tip [21]. In a first approach, the configuration can be described by a series of two springs, one representing the tip–sample contact stiffness kts, and the other representing the cantilever stiffness kc (Fig. 8.1 c). The tip–sample contact stiffness is the one which is to be characterized. When the sample is displaced by a certain distance Dz in the vertical direction, both the tip–sample spring and the cantilever spring are compressed. The deformation of the cantilever spring is detected. It is a measure of the tip–sample contact stiffness, since the higher the latter, the higher the deformation of the cantilever. Similarly, in a dynamic version, the vertical sample position is modulated sinusoidally and the resulting dynamic deformation of the cantilever is evaluated. Usually, the lock-in technique is employed for signal recovery since it allows a highly effective frequency-selective amplification and output signals are provided which are proportional to the amplitude and to the phase shift, respectively. Thus, images of the amplitude and phase of dynamic cantilever bending can be recorded simultaneously with the topography image. In general, this dynamic technique of nano-indentation is referred to as force modulation microscopy (FMM) [22, 23]. Strictly speaking, the displacement modulation approach elucidated above must be distinguished from other approaches where direct force modulation is applied via an oscillating magnetic field [24, 25]. The FMM results given in the following section were acquired using the displacement modulation approach. When soft materials such as polymers are being investigated, it provides sufficient sensitivity and there is no requirement to use magnetized cantilevers. 8.2.2 Depth-Sensing Micro-indentation (DSI)

In order to determine the local mechanical properties such as the hardness and modulus of non-flat samples or heterogeneous multiphase materials, a DSI system was integrated in a scanning device. This system combines SFM-like topography imaging with the ability of DSI tests to be performed at selected areas of interest by using a well-defined diamond tip. Basically, a force transducer (TriboIndenter; Hysitron Inc., Minneapolis, MN) was mounted on the fixed outer frame of a piezoelectrically driven xy-scanning stage (max. displacement 175 lm) with a large opening (Fig. 8.2). The sample is

8.2 Stiffness Mapping by Indentation Techniques

Fig. 8.2 Schematic representation of the integrated scanning– indentation system.

situated underneath the indentation tip and force transducer, mounted on top of a z-piezoelectric actuator (max. displacement 24 lm), which is used for height control of the sample. All the elements are controlled by capacitive sensors, and hysteresis is suppressed by closed-loop control. The force transducer can be adjusted in height and angle by three precision screws. For positioning of the force transducer, an optical microscope is mounted on the same frame as the force transducer. The microscope can be adjusted so that the central viewpoint is focused on the same position as the indentation tip. For imaging the topography of the sample surface, a multitask digital signal processor is used (DSP, Adwin Gold type; Jäger GmbH, Lorsch, Germany). When scanning, the DSP controls the xy-movement according to the user-defined scan field as well as the z-position of the sample in order to hold the force setpoint constant. The DSP is a multiprocess stand-alone system. It receives user-defined parameters from a Windows®-based host PC and sends back force data and the x-, y-, and z-positions. Homemade software is employed for scan operation as well as display and storage of the data. After the surface has been scanned, positions for DSI tests can be defined with the assistance of the topography information. For DSI tests, the xyz-stage acts as a positioning system. The site of the sample surface to be indented is moved underneath the tip, and then the sample is lifted to make light contact (load 2–6 lN). Afterward, motion along all axes is stopped and the standard indentation procedure of the Hysitron software package can take over control. When the DSI procedure is complete, the indenter is moved to the next position. Finally, when the complete list

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of test positions is processed, the indentation field can be scanned again (load 2–6 lN) in order to obtain the indentation imprint images. Examples are given in Figs. 8.5 a and 8.8 a below. The shape of the indenter has Berkovich geometry, a three-sided pyramid with an opening angle of 1368. In the ideal case, the projected area-to-depth function is given by the square of the depth of contact [26]. However, in general deviations from this simple relationship have to be taken into account, due to blunting of the apex of the tip. The particular area-to-depth function of the indenter being used was determined by means of an epoxy sample possessing a Young’s modulus of *3.08 GPa. Although the indenter is relatively flat, reasonable-quality images of the surface of a sample can be obtained.

8.3 Some Fundamental Aspects of Interphase Mapping by Indentation Techniques 8.3.1 Artifacts Induced by Topography

It should be made clear that the FMM signal can also be affected by factors other than the local sample stiffness. As is obvious from the equation for the tip–sample contact stiffness [Eq. (2)], variations of the contact radius a as well as variations of Ets will change kts. Changes in the tip–sample contact area can be caused by the sample topography both in a direct and in an indirect sense. In the direct sense, the geometry of topographic features such as hills or pits causes the contact area either to shift on the surface of the apex or to change in magnitude. Hills are characterized by a convex curvature and pits by a concave one. For instance, in pits with a radius of curvature comparable with that of the apex, the contact area will be significantly increased. Via the contact radius, this will result in an increased contact stiffness [see Eq. (2)]. In the indirect sense, the coupling occurs via the topography feedback. Owing to the nonzero period of time, the feedback needs to adjust the cantilever height position to the topography at the position of the tip; the normal force exerted by the tip on the surface is not perfectly constant but exhibits some deviations from the setpoint value. In the case of the tip scanning uphill, the noninstantaneous reaction of the topography feedback will lead to a somewhat increased normal force which in turn results in an increased contact area. However, this second effect via the topography feedback depends on quite a few parameters, such as the scan velocity and the various control parameters for adjusting the performance of the feedback. Therefore, these topography-related effects are hard to take into account in a quantitative manner, and serious data correction procedures will be difficult to achieve. Thus a more pragmatic but nevertheless rigorous approach was devised for ruling out topography-related effects on the FMM stiffness signal and the potential misinterpretation of stiffness data [18, 28]. The region of analysis (ROA) is limited to regions where the amount of the local surface angle of slope (as measured in the direction of scanning) is lower than a critical value. This

8.3 Some Fundamental Aspects of Interphase Mapping by Indentation Techniques

approach is motivated by the general finding that the image of normal forces resembles that of the local slopes of topography [18, 28]. That is, when scanning uphill, the surface slope angle is positive and the normal force is greater than the force setpoint value; similarly when scanning downhill. Moreover, the lateral forces also are affected by topography [27]. Analyzing a number of topography images recorded with a typical feedback performance, it was found that the maximum surface slope angles to which the feedback is able to adjust with nearly perfect accuracy are *|3|8 [18, 28]. The images were recorded with a typical scan frequency of 1 line s–1 and the distance over which the slope angle a was calculated *1/50 of the scan range. According to a = arctan(dz/dx), a was calculated from the slope which results by differentiating with respect to position along the scan line. Although better performance of the topography feedback may be achieved with peculiar SFM systems, a critical slope angle of *|3|8 can be defined and used in the sense of a conservative quality criterion. Considering that low value, it seems that with this criterion the direct contact area effects are also negligible, at least in the case of typical SFM tips with radius of curvature R less than *50 nm. It should be noted that the image of slope angle has to be calculated from the raw topography image as measured, i.e., without prior application of any software-based leveling techniques. The concept is confirmed by the following example of a wave-like structure on the surface of an epoxy sample (Fig. 8.3). For convenience the two regions of particular interest are labeled A and B, respectively. In region A the change of height is not as dramatic as in region B. As usual, great brightness indicates high altitude, and vice versa. Interestingly, the FMM image shows that the amplitude did not change in region A, but there was a change of amplitude in region B. Because the amplitude is homogeneous in the rest of the image, including region A except for the upper part of the image, where there are many grain-like structures, it seems certain that the amplitude variation in region B originates from the topo-

Fig. 8.3 Effect of topography on FMM amplitude, as measured on epoxy. (a) Topography image; (b) image of slope angles as calculated from the topography; (c) FMM amplitude image. In (b), two regions A and B are marked. In region B, where the slope angles

exceed +38 or –38, the amplitude image exhibits variations. Total height range of the topography image zmax = 97 nm; cantilever stiffness kc * 3 N m–1; FMM frequency 57.0 kHz; FMM amplitude * 0.1–1 nm.

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graphic features rather than from true stiffness variations. Figure 8.3 b shows the slopes on the surface, which were calculated from the topography image. After the calculation of slopes, a threshold value was sought, to find the critical slope above which topography-related artifacts occur in the FMM amplitude. As can be seen from Fig. 8.3 b, the critical slope value is *|3|8. This indicates that if the unevenness of the surface becomes greater than 38 or lower than –38, contrasts are induced in the FMM image. As a result, the FMM amplitude, which is supposed to contain stiffness-related information only, is also affected by topographic features. Some remarks should be made about the application of the topography criterion: · By its nature, the topography criterion defined above is necessary but not sufficient. The performance of the feedback at a certain point of time depends not only on the topography at the current position of the tip but also on the topography at the position it passed just beforehand. In terms of the widespread proportional-integral–differential (P-I–D) algorithm, the I and the D terms take account of the noninstantaneous reaction of the feedback. In particular, long-standing deviations from the force setpoint value affect the feedback via the I term. Correspondingly, at the edges of the ROA, significant deviations from the force setpoint value may occur due to topography features situated outside the ROA but still affecting the feedback [28]. · It should be borne in mind that the true topography differs from the measured one. Aside from the imperfect topography feedback, differences can result from the limited sharpness of the tip or from the nonzero compliance of the surface material. Whereas the latter two effects will not be reflected by the force image, it proves useful for checking deviations from the feedback setpoint value. · Apart from the normal force signal, the lateral force signal should also be inspected for topography-related force variations. Topographic slopes induce lateral forces, thus contributing to the torsional deformations detected in lateral force mode [27]. The related tilt of the tip means changes in the tip–sample contact area, in particular its position on the apex of the tip. If there is severe misalignment of the optical detection scheme, torsion cantilever deformation may even make a contribution to the normal force signal. Due to a limited number of data channels, however, in many cases the simultaneous measurement of FMM amplitude and phase, as well as normal and lateral forces, is not possible. Moreover, careful consideration of all these data can turn the data evaluation procedure into a time-consuming task. Hence, application of the topography criterion seems to be a valuable approach, in particular since contact area variations in the direct sense are not reflected by force images. The global tilt of the sample surface can easily be higher than the critical slope angle of *38 defined above, thus rendering impossible any stiffness analysis with strict application of the topography criterion. However, parallelization


Fig. 8.4 (a) The steps of sample preparation necessary for the EBL-based approach. The optical microscope images demonstrate the structures received after distinct steps. (b) SFM topography image of the region around a Cu/epoxy interface. The area of the Cu microstructure is located in the upper

part of the image. The Cu surface is * 20 nm higher than the epoxy one, which indicates that the epoxy formulation was not in contact with the Si substrate. Total height range of the image zmax = 56 nm; cantilever stiffness kc * 3 N m–1.

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between the plane of scanning and the plane of the sample surface can be achieved when using a setup where the cantilever is scanned rather than the sample. In the less favorable case where the sample is scanned, these two planes are not decoupled from each other and any tilt between the plane of the sample surface and the head of the microscope will result only in a tilt of the tip. It should be noted that in the case of pointwise stiffness measurement (i.e., without any simultaneous scanning motion), effects resulting from the sluggishness of the topography feedback are avoided. However, the contact area is still affected by the tilt of the sample surface. The Cu/epoxy sample used for the above analysis was prepared in a novel way, tailored to get particularly flat and easy-to-scan surfaces. In this recent approach, electron beam lithography (EBL) and the lift-off technique were employed in order to produce well-defined Cu structures [29]. The sample preparation procedure is illustrated in Fig. 8.4 a. For applying the EBL technique, a poly(methyl methacrylate) (PMMA) positive photo-resist (Allresist GmbH, Strausberg, Germany) was spin-coated on a Si wafer. The film had a thickness of *1.5 lm and it was pre-baked at 170 8C for 10 min. Once the substrate with the photo-resist was prepared, it was placed in a scanning electron microscope and exposed to the electron beam to write the patterns. After the exposure the substrate was developed to remove the exposed area, then Cu was deposited in a layer *500 nm thick by physical vapor deposition. By the lift-off technique the rest of the photo-resist was removed and the desired Cu patterns were obtained. The DGEBA-type epoxy resin (Epikote 828; Shell Chemicals Europe) and the curing agent (aliphatic and cycloaliphatic amines, Epikure F-205; Shell Chemicals Europe) were mixed well at 60 8C in a stoichiometric ratio (100 : 57 w/w). Then the mixture was carefully placed on top of the Si substrate with Cu structures. The epoxy system was cured at 80 8C for 2 h and the post-curing was done at 130 8C for 1 h. After the curing procedure, the Si substrate was removed by notching and shearing, and the resulting Cu/ epoxy surface was ready for analysis using SFM. As a flat substrate was used to obtain a flat surface, this approach is referred to in the following as the replica technique. 8.3.2 Artifacts Induced by the Extent of the Stress Field Beneath the Indenter

When small structures or heterogeneous materials are being tested, the problem often arises of how close to the phase boundary one phase can be tested without significant interaction with the neighboring phase. If the stress field beneath the indenter is significantly influenced by the neighboring phase, an apparent change in hardness and modulus values is observed which is merely a mechanical artifact. This artifact can easily be confused with real changes in the local properties of the materials, due to segregation effects, hindered mobility, transcrystallization, etc. In order to assess this kind of mechanical interaction, a two-phase sample of steel and polycarbonate (PC) was tested. It was assumed


that if the amorphous and chemically stable PC forms a mechanical IP, it is not wide enough to play a role in the study. The only interface effect to be expected is a thin region of altered chain conformations, with a width given by the average radius of gyration, i.e., some tens of nanometers at maximum. As compared with the micron-scale indentation marks discussed below, this seems to be negligible. A droplet of commercially available PC (Makrolon 2400c; Bayer AG, Ludwigshafen, Germany) was melted on a polished steel surface (St34 type) at 300 8C. After cooling, both parts were embedded into a common epoxy resin for further metallographic processing. After curing, the sample was cut perpendicularly to the steel surface using a diamond saw. It was polished using SiC emery paper of grades 1200 and 2400. However, the finish was done using a diamond lapping film (0.5 lm grade; 3M GmbH, Neuss, Germany) for *15 s. This preparation step ensures an in-plane steel/polymer interface region without severe height differences between the two components. After adjustment of its surface normal to the indentation axis, the sample was fixed to a magnetic holder. In order to find the interface region, an optical microscope was focused to the point where the indenter contacted the sample surface. This calibration was done by means of a comparatively large indent (maximum load 10 mN) in scratch-free aluminum foil. After the indentation, the microscope was adjusted to focus the visible triangular imprint. Then the area of interest was moved into the center view of the optical microscope. After the transducer had been replaced in the same position as before, the scanning process was started. Using the resulting topography images, suitable interface regions were selected and multiple series of indentations were performed in the PC phase. Within each series, indents were located at different distances from the steel/PC interface. After every series, the tested area was scanned again, and the diameter D of each individual imprint as well as the distance X between its center and the steel interface were determined. A typical situation after a test series is shown in Fig. 8.5 a. The indentation load–time regime was as follows: 5 s loading up to the maximum load, 30 s holding period, and 5 s unloading. Maximum loads were 0.8, 1.6, 3.2, and 6.4 mN, respectively. Each indentation was evaluated according to the procedure by Oliver and Pharr [26], which is included in the Hysitron software package. As shown in Fig. 8.5 b, the elastic response is a function of the distance from the phase boundary. For all maximum loads Fmax, the apparent increase in the stiffness becomes significant where the ratio X/D is less than unity. Within the tested range of forces, the critical distance with X/D = 1 increases linearly with Fmax. The lower Fmax, the closer to the interface the indentations can be performed without disturbing influences. In the case of Fmax = 0.8 mN, the critical distance is *8 lm, but when 6.4 mN is applied, the critical distance is *17 lm. Thus, as a rule of thumb, at least one imprint diameter should lie between the interface and the center of the indenter. These results can be interpreted in terms of the subsurface stress distribution beneath the indenter [30]. For instance, the vertical extent of the stress field be-

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Fig. 8.5 (a) Topography image of micro-indentations at an interface between steel (left) and PC (right). The indentations were made on PC along a line which was nonparallel to the interfacial borderline. (b) Plots of the measured values of the reduced Young’s

modulus Er versus the distance from the interfacial borderline. The width of the zone where enhanced modulus values are measured increases with the maximum load of the DSI experiment.

comes obvious from the fact that at indentation depths larger than *1/10 of the thickness of a coating, the measurement is influenced by the mechanical properties of the substrate [31]. The measurement can be influenced by a horizontal boundary; but it can also be influenced by a vertical boundary, as long as the subsurface stresses exceed a critical value at the position of the boundary.

8.4 Two Cases of Mapped Epoxy Interphases 8.4.1 The Cu/Epoxy Interphase

Two examples of the characterization of IP stiffness profiles using FMM are given in Fig. 8.6. In both cases mixtures of epoxy resin with aliphatic/cycloaliphatic amines were cured in the presence of a Cu component. The investigated surface was prepared either by a replica technique similar to that described in Fig. 8.4, or by sectioning of a bulk sample (see Fig. 8.6 a, b and Fig. 8.6 c, d, respectively). In a similar manner to the preparation technique described in Section 8.3.1, a Cu/epoxy composite sample was prepared from a Cu film deposited on a substrate. However, in this case Cu edges were generated by etching tiny holes into the Cu film by spraying droplets of a FeCl3 solution onto it. The substrate material was mica. The epoxy system was the same as in Fig. 8.3, but mixing and curing were done at 50 8C. Details of the sample preparation procedure are described in Ref. [17]. A corresponding FMM amplitude image is shown in Fig. 8.6 a. The stiff Cu surface is located in the left part of the image, as may be

8.4 Two Cases of Mapped Epoxy Interphases

Fig. 8.6 Two examples of epoxy IPs within Cu/epoxy composites. (a) and (c) The FMM amplitude images measured in the region of the Cu/epoxy interfaces. In both cases the Cu region is located in the left part of the images (high stiffness). (b) and (d) Results

of the fits in a pseudo-3D representation. The ROAs are marked with (a) black and (c) white borderlines, respectively. N denotes the direction perpendicular to the mean local tangent (T) of the interfacial border. Adapted from Ref. [18].

seen from the high amplitude. The gray scale was adapted to the range of amplitude values measured on epoxy. Near the interfacial borderline, a lower amplitude is observed on the epoxy than in the region further away (hereafter referred to as bulk epoxy). Within the ROA defined by the 3 8 topography criterion, the amplitude profiles were analyzed in the direction perpendicular to the mean local tangent of the borderline, i.e., along N. The results of the fitting procedure using a semi-Gaussian relationship are given in Fig. 8.6 b. In particular, the average total width of the amplitude gradient was deduced to be *280 nm. Another demonstration of the application of the topography criterion is given in the next example of the FMM analysis of a Cu/epoxy IP (Fig. 8.6 c, d). In that case the sample surface was prepared by cross-sectioning a Cu/epoxy composite using a diamond saw. The epoxy system was L180/H181 (both from Scheufler Kunstharzprodukte GmbH, Stuttgart, Germany). According to the information given by the manufacturer, the epoxy equivalent weight of the resin L180 ranges between 160 and 180 g mol–1. The curing agent H181 is of low viscosity (*0.01–0.02 Pa s at 25 8C) and contains aliphatic as well as cycloaliphatic amines. A 100 : 25 (w/w) stoichiometric mixture was prepared. Both mixing and curing were done at 50 8C. After curing and cutting, some polishing was required in order to reduce the surface roughness. A sufficiently extended ROA was found where the 3 8 topography criterion was met. The ROA is marked with a white borderline. In addition, the Cu edge is marked with a black border. Within the ROA the final analysis was limited to the area marked with thick white borderlines. Similarly to the case of the Cu/epoxy IP shown in Fig. 8.6 a and b, the epoxy stiffness was observed to decrease when approaching the interface from the bulk epoxy. However, the width of the IP was much greater in this case. Quantitative analysis by fitting a semi-Gaussian profile to the data delivered a total width of the IP stiffness profile of *29.4 lm. Although these two IPs possess different widths, in both cases the epoxy stiffness was found to be reduced in the vicinity of the interfacial border between Cu(oxide) and epoxy. In the second case, the existence of a much wider IP can

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be considered as an indication of different kinetics, even if the driving force for IP formation was the same, such as preferential adsorption of the amine curing agent. The differences between the two samples are the epoxy systems (Epikote828/F205 versus L180/H181) and the techniques for preparation of the surfaces to be investigated (replica versus cutting). It should be noted that in the case of the replica sample, curing was performed in the presence of the mica; however, a gap of *110 nm existed between the epoxy and the mica surface, as is obvious from the corresponding topography image (not given here; see Ref. [17]). Hence interface-related effects due to the presence of mica can be widely ruled out. 8.4.2 The PVP/Epoxy Interphase

Another class of adherends is that of thermoplastic polymers. In contrast to metal adherends, thermoplastics are not impenetrable and thus absorption effects can be expected in addition to adsorption phenomena. Hence, given sufficient conditions for preferential absorption, a considerable mass uptake by the thermoplastic can occur, potentially resulting in significant stoichiometric imbalances on the epoxy side. Apart from the driving force for absorption of molecules from the liquid epoxy formulation, it is the diffusivity of these molecules within the thermoplastic which plays a major role in the interdiffusion process. In particular, the diffusivity is affected by the mobility of the host molecules. Thus enhancement of diffusivity occurs in the glass transition region and at higher temperatures when intermolecular cooperative motion is activated. As a case study, a liquid epoxy formulation was cured in the presence of a film of the thermoplastic polyvinylpyrrolidone (PVP). The epoxy system was a stoichiometric mixture of diglycidyl ether of bisphenol A (DGEBA) with diaminodiphenylsulfone (DDS). As shown by Oyama et al. by means of energy-dispersive analysis of X-rays (EDX), at the epoxy/PVP interface pronounced interdiffusion effects can be expected [7]. These were attributed to strong chemical interactions between epoxy and the carbonyl group of PVP. Owing to the S atom in DDS, concentration gradients of the amine curing agent are detectable by EDX, using the S atom as a tracer element for amine molecules. In this way, concentration gradients extending over several microns were found [7]. In the present study, the focus was on the IP modulus variations corresponding to the chemical IP. With the purpose of obtaining quite wide IPs, the epoxy curing was performed at a temperature in the region of the glass transition of PVP, i.e., at 170 8C [32]. The K-90 grade PVP used for preparing the films has a number-average molecular weight Mn of *360 000 [33]. As deduced from differential scanning calorimetry (DSC) measurements of the PVP–K90 film, the center temperature of the glass transition was at *169 8C, with the onset at *158 8C. Typically, the thickness of the PVP films was in the range 50–150 lm. Sandwich-like samples were prepared with the PVP film covered on both sides with epoxy. After curing, cross-sections in the direction perpendicular to the

8.4 Two Cases of Mapped Epoxy Interphases

film were prepared using an ultramicrotome equipped with a diamond knife. An SEM micrograph and corresponding concentration maps resulting from EDX are given in Fig. 8.7. With some effort, the edges of the PVP film are visible in the micrograph. As a visual aid, these edges are marked with white dashed lines, which are also drawn on the EDX concentration maps shown in Fig. 8.7 b and c. From the Nmap (Fig. 8.7 b), the PVP layer is the one exhibiting a higher concentration of N atoms, whereas the epoxy layers are identified as the ones with a high concentration of S atoms (Fig. 8.7 c). However, as is obvious from the S map, the concentration of S is strongly reduced within a comparatively wide zone adjacent to the PVP/epoxy interface. Analysis of profiles perpendicular to these interfaces reveals that the widths of these concentration gradients are *73 lm and *61 m for the IPs on the left and on the right, respectively. The variation of these IP widths may reflect differences between the upper and the lower side of the film surface, e.g., due to the fact that the film was made by casting the corresponding solution into a Teflon mold. Thus, one side of the film was in contact with Teflon, whereas the other one was in contact with air. As can be seen from the cross-sectional profiles resulting from averaging over the rectangular areas marked in Fig. 8.7 b and c, there is a nonzero concentration of S atoms

Fig. 8.7 (a) SEM micrograph of the crosssectional surface of an epoxy/PVP/epoxy sandwich sample. (b) EDX map of the local N concentration, indicating the borders between PVP (bright) and the epoxy layers. The borders are marked with white dashed lines. (c) EDX map of the local S concentration of the same region as in the SEM

micrograph. S is contained in the amine curing agent DDS. Obviously, close to the interfaces the S concentration is significantly lower than in bulk epoxy. The cross-sectional profiles given on the graph (N: black line; S: gray line) resulted from averaging over the rectangular area marked with a white dotted line.

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within the PVP layer. Conservation of mass requires that the amine molecules lacking in the epoxy layers were absorbed by the PVP layer, the width of which is *61 lm. In this case, owing to the rather wide extent of the amine concentration gradients, a micro-indenter was employed for mapping IP stiffness variations. A diamond tip with Berkovich geometry was used. On the one hand it is less sharp than an SFM tip, but on the other hand its shape is well defined and the calculation of modulus values is possible within acceptable margins of error. The Young’s modulus was deduced from the initial part of the retraction curve of the force–penetration experiments. At maximum load a hold period of 30 s was obeyed in order to separate creep from purely elastic recovery. The maximum indentation force was *10 mN. Since the total extent of the epoxy IP is greater than the maximum scan range of the DSI setup, the scan position was translated several times and successive scans were stitched together. The total width of the stitched scan field is 290 lm. The resulting topography image of the epoxy IP as measured after the DSI experiments is given in Fig. 8.8 a. On the surface of the PVP layer, the indentation marks are readily observed to be larger than on the epoxy surface, indicating a much lower hardness. A plot of the deduced sample modulus values, Es, versus the distance from the epoxy/PVP border is shown in Fig. 8.8 b. The coordinate xi–x denotes the distance from the interface border, located at (xi, yi) and marked in Fig. 8.8 a with a solid black line. Two interesting findings can be reported. First, a monotonic in-

Fig. 8.8 (a) Topography image of a field of micro-indentations across the left-hand epoxy IP. The interface is marked with a black solid line, on the right of which is located PVP. The experiments were performed using the setup described in Section 8.2.2. (b) Pseudo-3D representation of the changes in the epoxy Young’s modulus, Es, across the epoxy IP. The Es values (dots)

displayed were calculated from the measured values of the reduced modulus, Er, using Eq. (3) and the values Et = 1140 GPa, mt = 0.069 for the diamond tip. The plane resulted from the fit of an error function to the data. xi–x denotes the distance from the interfacial border, (xi, yi). The total modulus change across the epoxy IP is * 1.1 GPa.

8.5 Conclusion

crease in epoxy modulus is observed with decreasing distance, xi–x, from the border. Second, the total width of the modulus gradient is *175 lm, i.e., about 2.6 times more than that of the amine concentration gradient detected by EDX analysis. Thus, the IP stiffness variations are hard to explain merely by the offstoichiometry effect alone [32]. Checking of the length calibration of both the scanning electron microscope and the scanner employed for the DSI measurements showed that the different widths of the chemical and the mechanical IPs cannot be traced back to miscalibrations. Finally, it should be noted that the diameter of the indentation imprints was much smaller than the width of the IP, and sufficient lateral resolution could be achieved by means of the DSI technique. In order to elaborate the relationship between stoichiometric imbalance of the epoxy–amine system and its resulting elastic properties, a series of reference samples of well-defined concentration ratios r was investigated using dynamic mechanical analysis (DMA). A clear increase in the glassy epoxy modulus (as measured at 20 8C) with an increasing excess of epoxy (r < 1) was observed [32]. Notably, the observation of modulus enhancement with increasing excess of epoxy is consistent with the concentration and modulus maps of the PVP/epoxy IP.

8.5 Conclusion

Interphase stiffness profiles were measured on epoxy. Basically, two different systems were investigated, namely Cu(oxide)/epoxy and PVP/epoxy joints. As a curing agent, aliphatic and cycloaliphatic amines were used in the first case and an aromatic one in the second. The epoxy stiffness was found to be reduced within the interphase in all the cases of Cu/epoxy interfaces analyzed. In contrast, in the case of the PVP/epoxy system, the stiffness of the interphase epoxy was found to be strongly enhanced. Across the interphase, the total change in the Young’s modulus was *1.1 GPa. From the EDX and DSI results it can be inferred that the interphase formation was driven by the preferential absorption of amine molecules and that strong spatial variations of local mechanical properties can result from the chemical gradient. In future, the calculation of interfacial stress distributions in epoxy matrix composites can benefit from such data. For mapping the interphase stiffness profile with sufficient lateral resolution, quite a number of indentations is necessary. A higher lateral resolution is required for narrow interphases and strong stiffness gradients. However, indentations with large penetration depths generate more widely extending subsurface stress fields. As a consequence, it is necessary to comply with larger minimum distances from the interfacial borderline. However, indentation experiments beyond the elastic regime provide access to further information such as hardness. Care has to be taken not to confuse apparent stiffness variations induced by topographic features with real stiffness gradients. In a rigorous approach, it was

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proposed to constrict the region of stiffness analysis to areas where the topographic slopes are limited.

Acknowledgments

We express our gratitude to S. Megard (formerly BAM, VI.21) for the performance of the DSI measurements at the steel/PC interface. The EDX data were partly collected by G. Eltanany (formerly BAM, VI.21) in the framework of her Ph.D. studies, which she did not continue, however. We gratefully acknowledge her contribution. Furthermore, we express our gratitude to D. Neubert (BAM, VI.11) for running the DSC experiments. The following chemicals were kindly provided: the Epikote828/Epikure F205 epoxy system by Brenntag GmbH; and the DER332 epoxy resin by Nordmann–Rassmann GmbH.

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31 32 33

Polymere, Verlag für neue Wissenschaft, Bremerhaven, Germany, 2002, p. 52. J. Chung, M. Munz, H. Sturm, submitted. B. Bhushan, Principles and Applications of Tribology, Wiley-Interscience, New York, NY, 1999, p. 198 ff. N. G. Chechenin, J. Bottiger, J. P. Krog, Thin Solid Films 1997, 304, 70–77. M. Munz, H. Sturm, W. Stark, accepted for publication in Polymer. V. Bühler, Kollidon – Polyvinylpyrrolidone for the Pharmaceutical Industry, 6th edn., BASF, Pharma Ingredients, Ludwigshafen, Germany, 2001, p. 36.

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9 Mechanical Interphases in Epoxies as seen by Nondestructive High-Performance Brillouin Microscopy J.K. Krüger, U. Müller, R. Bactavatchalou, D. Liebschner, M. Sander, W. Possart, C. Wehlack, J. Baller, and D. Rouxel

Abstract

Boundaries between a reactive polymer system, e.g. an epoxy, and solid substrates modify the mechanical properties of the polymer in the interface-near region. The so created interphases are of unexpected width up to several hundred lm. The recently developed Brillouin microscopy grants nondestructive access to the local mechanical properties within these interphases. It turns out that these mechanical interphases depend strongly on the kind of polymer system, the substrate material and the preparation conditions.

9.1 Introduction

Polymer networks such as epoxies play an increasing role as adhesives in industry. Two properties are of special importance for their application: (a) a strong adhesive bond is required between the solidified adhesive and the bonded object, which is often a metal; (b) the mechanical stiffness of the adhesive has to be adapted to the desired level. As a consequence, the adhesive has to be selected according to its adhesion properties as well as its mechanical properties. Several studies have shown that both properties are linked as soon as the epoxy polymer layer is sufficiently thin: the contact of the polymer with the substrate may induce in the polymer a broad interphase where the morphology is different from the bulk. Roche et al. indirectly deduced such interphases, for example from the dependence of the glass transition temperature on the thickness of the polymer bonded to a metal substrate [1]. Moreover, secondary-ion mass spectroscopy or Auger spectroscopy provided depth profiles of interphases in terms of chemical composition, which showed chemical variations at up to 1 lm distance from the substrate. In this paper, Brillouin microscopy (BM), an original technique which provides information about spatial variations of mechanical properties, is preAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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9 Mechanical Interphases in Epoxies as seen by Nondestructive High-Performance Brillouin Microscopy

sented. The power of this new technique is shown in providing outstanding information about mechanical interphases with a spatial resolution down to about 1 lm. A room temperature (RT)-curing amine–epoxy system (mass ratio DGEBA/DETA = 100 : 14) is used as a model adhesive. The influence of different metal substrates on the formation of mechanical interphases in epoxy polymers is demonstrated and discussed.

9.2 Brillouin Spectroscopy on Thermal Phonons and Other Elementary Excitations 9.2.1 An Introduction to the Physics of Classical Brillouin Spectroscopy

Brillouin spectroscopy (BS) is an optical technique which is used predominantly to investigate acoustic properties at hypersonic frequencies. BS can be applied only for transparent materials, at best for translucent materials. The acoustic wavelengths involved range from about 200 nm to several microns. Fig. 9.1 shows a typical Brillouin set-up. For one of a number of possible optical arrangements (cf. Fig. 3.2), the laser illuminates the sample. The scattered light is collected by suitable optics and the spectral intensity distribution of the inelastically scattered part of the light is analyzed with a Fabry-Pérot spectrometer. The optoelectronic conversion is performed with a photomultiplier and the spectrum is accumulated on a multichannel analyzer (MCA) (Fig. 9.1). On the screen of the MCA, a schematic drawing of a typical Brillouin spectrum is shown. Generally the intensity I* q; x of scattered light per solid angle and frequency interval is proportional to the space and time Fourier transform Fr;t fAg of the autocorrelation function A of the optical polarizability fluctuations dakl * r; t [2] [Eq. (1)], as is expressed by Eq. (2). r; t damn * r 0 ; t0 i Aklmn hdakl *

1

Is;i * q; x / Fr;t fAg

2

The subscripts s and i denote scattered and incident light, respectively, and refer to the appropriate optical polarization directions of the light waves here [3]. Omitting, for simplicity, the tensor properties of da we see the spectral power density Is;i * q; x to be proportional to the mean square fluctuation component da* q at frequency x [Eq. (3)]. Is;i * q; x / da* r 2x

3

All kinds of excitations, e.g., phonons, excitons, spin waves, may contribute to da* q. Higher-order processes such as multi-phonon interactions may be in-

9.2 Brillouin Spectroscopy on Thermal Phonons and Other Elementary Excitations

Fig. 9.1 Schematic representation of a Brillouin spectrometer for measurements in three different scattering geometries: backscattering (1808), 90R scattering and 90A scattering (beam path defined by shutters and mirrors). The laser power can be modu-

lated by a Pockels cell; the sample holder provides a defined temperature (T) as well as rotational (a) and translation (Dx) scans. Additionally, scattered light may be analyzed by a polarizer.

cluded formally in this treatment by expanding da into a power series in terms of symmetry coordinates. Assuming a certain elementary excitation characterized by an extensive parameter w* q; x, a conjugated force F* q; x, and a susceptibility v* q; x [Eq. * (4)], the spectral power density I q; x can be related to the imaginary part of v* q; x by the fluctuation-dissipation theorem stated in Eqs. (5) and (6) [4]. w* q; x v* q; x F* q; x

4

I* q; x / nx 1 Imv* q; x

5

where 1 x h exp nx kT

1

6

The relationship may be extended to include several, say s, coupled modes wj . To obtain the resulting field of generalized forces, one has to add the contributions of these s modes [Eq. (7)]. Fi

s X j1

cij wj

7

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128


Here the diagonal elements of the matrix fcij g represent the inverse susceptibilities of the uncoupled modes. The off-diagonal elements characterize the complex mode–mode coupling strength. Inverting Eq. (7) one obtains the mode amplitudes [Eq. (8)]. wk

s X

vki Fi

8

i1

The susceptibilities vki determine the scattered light spectrum [Eq. (9)]. " *

I q; x / nx 1 Im

s X

# *

pi pj vij q; x

9

i;j1

The coefficients pi are related to the light-scattering cross-sections of the various modes and can be described, for example, by appropriate components of the elasto-optic tensor. The specific form of the light-scattering spectrum will depend mainly on the characteristics of the inverse susceptibilities cjj of the uncoupled modes. Taking Eq. (10) for a damped harmonic oscillator, and Eq. (11) for a relaxator, it is easy to predict the form of the scattering spectra. cjj cj0 x2j

x2 i x C j

10

cjj cj0 1 i x s

11

However, experiments do not give sufficient information for an unambiguous analysis. Even in the two-mode case, the complex factor c12 cannot be deduced definitely if coupling appears. From Eqs. (10) or (11) and (9), the spectral power density of an uncoupled relaxator or oscillator can be deduced. The imaginary part of the inverse of Eq. (10) is given by Eq. (12). This corresponds to the equations of a damped harmonic oscillator with the temporal attenuation parameter C i. Imvii Osc ci01

xC i x2i

x2 2 x2 C 2i

12

The imaginary part of the inverse of Eq. (11) is given by Eq. (13). Imvjj Rel Imcjj 1 cj01

xC R x2 C 2R

13

Combining the spectral components of a slightly damped harmonic oscillator [Eq. (12)] and a relaxator [Eq. (13)] with Eq. (9) yields the total spectral power density [Eq. (14)].


" *

I q; x / nx 1

cj01

p1 xC R p2 xC Osc ci01 2 x2 C 2R xOsc x2 2 x2 C 2Osc

# 14

Eq. (14) would give the shape of a spectrum typical for a low-viscosity liquid, consisting of a central peak due to entropy fluctuations and frequency-shifted Stokes and anti-Stokes lines related to density fluctuations resulting in a bulk phonon. 9.2.2 The Kinematic View of Brillouin Spectroscopy

The kinematic view of BS couples energy and momentum of the photons and phonons interacting in the scattering process. As usual for inelastic scattering, the energy and momentum conservation principles hold [Eqs. (15, 16)], where xs ; xi ; X are the angular frequencies of the scattered light, the incident laser * * q are the correlight, and the interacting sound wave, respectively, and ks ; ki ; * sponding wave vectors [5, 6]. hxs hxi hX

15

hks hki hq

16

In BS, the energy transfer between photons and phonons is very small * * (5 ´ 109 Hz/5 ´ 1014 Hz = 10–5) and hence ki and ks have the same length. Provided, acoustic attenuation is small (C Osc X the phase-sound velocity can be calculated from simple geometric arguments. Equations (15, 16) yield Eq. (17), where n is the refractive index of the sample, for the sound velocity v for the longitudinal mode. Equivalent equations hold for other polarizations. v

X X kLaser q 2p 2n sin h2

17

Knowing the mass density q of the sample, the longitudinal elastic modulus q, can be calculated according to Eq. (18). This c, related to the wave vector * equation is most important since it provides the key for the determination of the mechanical properties of materials at hypersonic frequencies from Brillouin spectra in a nondestructive manner. q c~ q q v2 *

18

9.2.3 Scattering Geometries and Other Pitfalls

As is seen from Eq. (17), the acoustic wavelength K for isotropic materials [Eq. (19)] depends on the vacuum laser wavelength kLaser, the scattering angle hi within the sample, and the refractive index n of the sample.

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K

2p kLaser q 2n sin h2i

19

As a consequence of the influence of the refractive index, the acoustic wavelength is not constant with respect to temperature, pressure, etc. This may cause serious problems in interpreting Brillouin data. Fortunately, there exists the 90A-scattering geometry (see Fig. 9.2 b) which leaves the acoustic wave vector independent of the refractive index (for isotropic samples). For the scattering geometries shown in Fig. 9.2, the relationships between the sound velocity, the sound frequency, and the sound wavelength are given by Eqs. (20, 21, 22) respectively, with sin(q1/2) and v180(T) defined in Eqs. (23, 24) [7].

v90N T

X90N T kLaser 2p 2nT sin h2i

20

v90A T

X90A T kLaser 2p 2 sin h2i

21

v90R T

X90R T kLaser 2p 2nT sin h2i

22

with hi sin cosfarcsinsin458g 2 v180 T

X180 T kLaser 2p 2nT

Fig. 9.2 Typical scattering geometries used for Brillouin spectroscopy. (a) 90N-scattering geometry; (b) 90A-scattering geometry; (c) a combination of the 90A- and the 90R-scattering geometry; (d) back scattering geometry.

23

24

Notation: indices i, s denote incident and * * scattered quantities; ki , ks : wave vectors of * the laser light; q: wave vector of the sound wave.


Except for Eq. (21), the acoustic wavelength K depends on the refractive index and hence on the measured sound frequency for all scattering geometries. In the case of Eq. (21), the refractive index enters the refraction process and the phase velocity in a compensating manner and therefore the sound wavelength becomes independent of n. Provided sound dispersion is absent, Eq. (21) can be used in combination with Eqs. (20), (22), or (24) to determine the refractive index from Brillouin data, e.g., as in Eqs. (25, 26).

nT

p X180 T 2 90A X T

25

nT

p X90N T 2 90A X T

26

In the case of film- or plate-like samples with reflecting substrates on one side, there is a scattering geometry which is an alternative to the geometries shown in Fig. 9.2 and which we call the RIhA-scattering geometry [8]. It combines the advantages of the 90A and the backscattering technique. A special version of the RIhA-scattering geometry is needed for Brillouin microscopy (BM) and will be called RIBM scattering. A schematic drawing of this RIBM scattering is shown in Fig. 9.3. Unfortunately, in this scattering geometry the acoustic wave vector depends on the refractive index again. The RIBM scattering involves * two scattering vectors ki , given by the incident and by the reflected laser beam. Within the sample the scattering wave vector * q is given by Eq. (27). * h qj 2 j ki j sin j* 2

27

The scattering angle h depends on the outer refraction angles ai, as and on the refractive index of the sample [Eq. (28)]. h p

sin ai arcsin n

sin as arcsin n

28

The longitudinal hypersonic velocity vl of an isotropic sample is calculated from Eq. (27), provided the outer angles ai and as are specified, according to Eq. (29) where fl is the sound frequency of the longitudinal sound mode.

vl

Xl kLaser fl sin p q 2n

sin ai arcsin n

sin as arcsin n

1

29

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132


Fig. 9.3 Generalized scattering geometry for Brillouin microscopy. Notation: indices i, s for incident and scattered quanti* * * wave vector of ties; ki , ks : wave vectors of the laser light; q: the sound wave; ai, as and bi, bs: outer and inner angles of refraction; h: light-scattering angle; gp: glass plate; n: the refractive index.

9.2.4 Brillouin Microscopy

Brillouin microscopy (BM) is a special version of acoustic scanning microscopy. A suitable realization has been presented recently [9]. As shown in Fig. 9.3, it consists of the classical entrance section, which enlarges the working distance and contains the glass plate (gp); this plate can be rotated around an axis in order to shift the scattering volume along the incident laser beam. It is followed in the light path by a customized optical magnifier (not shown), which magnifies the image of the scattering volume to the desired value and projects it onto the entrance of a high-performance tandem Brillouin spectrometer [7]. The relationship between the size of the pinhole at the spectrometer entrance and the optical magnification factor of the optical magnifier determines the spatial resolution of the microscope. The third section is a modified high-performance sixpass tandem Brillouin spectrometer of the Sandercock type [7]. This TBS is fully automated. The temperature control of the Fabry-Pérot spectrometer permits long-term measurements over weeks [10]. The minimum information volume utilized by our BM has a lateral dimension of about 1 lm. It should be stressed that such a high spatial resolution is at the cost of scattered intensity. Hence, weakly scattering samples may cause problems for BM but polymers and liquids usually possess sufficient scattering efficiency. As a first experimental example, Fig. 9.4 depicts the spatial distributions of sound velocity v and sound attenuation C across a blown film of poly(4-methyl1-pentene) (P4MP-1, thickness 180 lm). There are two maxima at the surfaces and two adjacent local minima within the film. Then, the sound velocity decreases from left to right inside the film, indicating a different thermal or mechanical history in different parts of the P4MP-1. The sound velocity maxima


Fig. 9.4 Sound velocity v (black squares) and sound attenuation C (open circles) of a P4MP1 blown film as a function of the spatial position d of the scattering volume inside the film (see text for further explanation).

are accompanied by an excess of C. This result is not in favor of a relaxation process but rather indicates an increased density of acoustic scattering centers close to the surfaces of the polymer film. These scattering centers could be microcrystallites at the surfaces of the film. In addition, we suppose that the asymmetry in v(d) inside the film is caused by temperature differences at the two film surfaces during the blow process. Fig. 9.5 provides a BM measurement for the diglycidyl ether of bisphenol A (DGEBA) in contact with the native Al2O3 surface of an Al substrate. DGEBA is a viscous liquid with a melting point Tm = 315 K and a density q = 1157 kg m–3. This material has a strong tendency to vitrify at a thermal glass transition temperature Tg of *243 K. Furthermore, DGEBA is known to show only very weak interactions with native Al surfaces [11]. Longitudinal sound velocity and attenuation are measured as a function of the distance d from the metal at ambient temperature (i.e. well below Tm). Near the interface (d < 250 lm), neither v nor C depends on the position of the scattering volume in the liquid. There is only some scatter of the data due to the small light-scattering cross-section of DGEBA. This is exactly what has to be expected for an isotropic liquid. Any interaction with the solid will not induce a measurable long-range effect on the structure or on the density of the liquid. Hence, this simple experiment proves that our BM results are reliable even at the borders of a sample.

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Fig. 9.5 Liquid monomer DGEBA layer (1 mm thick) on an Al surface (at d = 0 lm): sound velocity v (black squares) and sound attenuation C (open circles).

9.3 Mechanical Interphases at Polymer–Substrate Interfaces 9.3.1 The Polymer Model System

The epoxy (EP) adhesive under investigation is composed of the resin DGEBA (D.E.R. 331; Dow) and the hardener DETA (diethylenetriamine; Fluka) with a mass ratio of 100 : 14. The oxirane rings of the DGEBA react with the primary and secondary amine groups of the DETA at room temperature (RT). As each of the components has several reactive groups, a three-dimensional network is formed in the course of the chemical reaction. A reasonable reproducibility is achieved if strict preparation instructions are obeyed. First, the DGEBA is heated to 55 8C and stirred for 30 min. It is also degassed by evacuation. Then the hardener DETA is added and the reactive mixture is stirred at 55 8C for 3 min in a N2 atmosphere to exclude humidity. This mixture is ready for application and curing at ambient temperature. This yields a limited chemical conversion of about 70% of the oxirane rings because the system undergoes a chemical glass transition which then drastically slows down the reaction. A post-curing step is applied for 1 h at 120 8C in a nitrogen atmosphere in order to achieve a fully crosslinked state. 9.3.2 Epoxy/Silicone Rubber Interphase

Simple tests show that the EP does not adhere to the silicone rubber Silgel (Elastosil 604; Wacker Silicone, Germany). However, DETA easily penetrates into the Silgel. This will influence the DETA exchange between the silicone and the EP sys-

9.3 Mechanical Interphases at Polymer–Substrate Interfaces

tem during curing. To study the effect on the state of the spatial network within the EP, we cured samples inside a Silgel cuvette. Fig. 9.6 shows the sound velocity profile for an epoxy plate (thickness 600 lm) cured and post-cured in a horizontal Silgel cuvette (wall thickness 1 mm). According to Eq. (18), v is related to the mechanical modulus of the network, which is known to increase with the density of crosslinks. The profile resembles an asymmetric W. The two minima, each at a distance of 70–80 lm from the EP–silicone interface, result from coupled processes. First, DETA diffuses from the monomer mixture into the Silgel while the cuvette is being filled and during the initial stage of curing. Hence some of the interphase region is depleted of amine and the density of crosslinks is reduced in the growing EP network. Simultaneously, the chemical curing reactions consume amine and this change of concentration induces a reflux of DETA from the Silgel into the EP sample. As the result, the crosslinking rises in the vicinity of the EP–silicone interfaces, thus providing the two maxima. Additionally, the vitrification of the EP network slows down the back-diffusion and the consumption of DETA before the minima in v(d) are filled up. The plateau region between the two minima is interpreted as the undisturbed bulk state of the epoxy sample since this region is not affected by the flux and reflux of DETA. This interpretation is confirmed by the data given in Fig. 9.7. It shows the temporal evolution of the elastic properties at three points in the EP sample during the curing process at room temperature. For that purpose we put a Silgel cuvette in an upright position, filled it with the DGEBA–DETA mixture, and started the BM measurement immediately afterward. The scattering volume was periodically positioned in the DGEBA–DETA mixture at different points: first at the Silgel–epoxy interface, second at 330 lm, and third at a 660 lm distance from the interface (see the insert in Fig. 9.7). It is likely that the temporal and spatial evolution of the sound velocity reflects the temporal and spatial DETA distributions within the sample. At tc = 0, the sound velocity values are almost the same for all positions of the scattering volume, but they start to deviate from each other with time. The crosslinking appears to be fastest at

Fig. 9.6 Profile of sound velocity v as a function of position d inside an epoxy plate cured and post-cured in a Silgel cuvette. The left and right sides of the graph correspond to the top and bottom of the horizontal cuvette, respectively.

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136


Fig. 9.7 Sound velocity v as a function of curing time tc during the crosslinking reaction of epoxy in a horizontal Silgel cuvette at room temperature, measured at three distances d from the Silgel–epoxy interface. The insert relates the three positions of the scattering volume to the corresponding profile v(d) given in Fig. 9.6.

d = 660 lm, as the slope of v (600 lm) at the beginning is the steepest and v itself reaches the highest level. The other two curves behave quite similarly at the beginning. Clear deviations appear after about 10 000 s. At d = 330 lm, v reaches a level which is 1% less than at the interface. At d = 660 lm, the sound velocity is 0.6% above the value at the interface. The temporal and spatial distributions of DETA in the EP sample depend on the geometry of the sample and of the cuvette as well. These geometrical dimensions define not only the storage properties for DETA but also the maximum geometrical paths for the DETA flux and reflux. This is illustrated in Fig. 9.8. Compared with the measurements shown in Fig. 9.6, the Silgel cuvette is now in an upright position, the wall thickness has been increased by a factor of 3 and the EP thickness by a factor of 7.5 for the data given in Fig. 9.8. In addition, this sample was not post-cured and hence the bulk value of the sound velocity is about 5% lower than in Fig. 9.6. The left minimum (0.8% below the plateau level) is found at 960 lm, which is an astonishingly large distance from the left interface. This result demonstrates the dramatically long-range influence of DETA diffusion on the elastic properties of the EP. 9.3.3 Epoxy/Metal Interphases

Mechanical interphases at polymer–EP interfaces are of great academic interest. Additionally, they are of technical importance in the case of metal–EP interfaces. Therefore, the formation of interphases was studied in different native metal–EP systems.

9.3 Mechanical Interphases at Polymer–Substrate Interfaces Fig. 9.8 Spatial sound velocity profile v(d) of an epoxy plate 4.5 mm thick, cured in an upright position in a Silgel cuvette at ambient temperature. The wall thickness of this cuvette was about 3 mm. For further explanation see text.

9.3.3.1 Technical Bulk Metals: Cu, Al The set-up for the epoxy sample preparation is illustrated in Fig. 9.9. Technical alloys of copper or aluminum plates (bulk) were polished to optical grade and used as one cuvette cover. The other cover was made of a PTFE plate with a polypropylene (PP) film as an interlayer in order to improve the optical properties of the adjacent EP surface. EP (DGEBA : DETA 100 : 14) was cured in the cuvette at ambient temperature. After removal of the PTFE plate and the PP film but still at ambient temperature, the 1 mm EP layers were subsequently characterized with BM by shifting the scattering volume across their thickness. As illustrated by the BM data in Fig. 9.10 for two independently prepared Cu– EP samples, the reproducibility of the measurements is excellent. According to Fig. 9.10, no mechanical interphases are observed at the polished Al and Cu alloys, neither in the RT-cured state nor after subsequent post-curing. Post-curing just increases the sound velocity. The adhesion strength of these EP–metal bonds is very low, however, and the related metal surfaces show no traces of EP and still look very well polished after fracture. Apparently, there is no chemical adhesion of EP to the polished copper and aluminum substrates.

Fig. 9.9 Cuvette for the preparation of EP–metal samples consisting of a PTFE plate, covered with a polypropylene film to assure optical quality, a PTFE spacer ring (d = 1 mm), and a polished metal plate (Al or Cu) as the substrate for the epoxy.

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138

9 Mechanical Interphases in Epoxies as seen by Nondestructive High-Performance Brillouin Microscopy Fig. 9.10 Sound velocity profiles v(d) for the RT-cured and for the post-cured state of two independently prepared EP plates on polished Cu alloy (open squares and triangles) as well as for an EP plate cured on polished Al alloy at ambient temperature (black squares). The metal is positioned at d = 0 mm.

For the RT-cured Al–EP sample, the v data are significantly lower than for the corresponding state on the copper substrate. Although an interpretation cannot be given yet, that difference indicates a homogeneous influence of the substrates on the whole EP layer.

9.3.3.2 Thin Evaporated Metal Substrates: Al, Cu, Au, Mg In comparison with native bulk copper and bulk aluminum substrates we have also prepared layers 100 nm thick of pure aluminum, copper, gold, or magnesium evaporated on silicon wafers. Due to the contact with ambient air, the native layers of oxides and adsorbates form on top of the metals. On these metal film substrates, EP plates 1 mm thick were prepared by curing the DGEBA– DETA mixture (100 : 14) at ambient temperature. As depicted in Figs. 9.11 and 9.12, the BM investigations reveal broad, mechanically stiffened interphases in the EP adjacent to all the metal films. The v level of these interphases is almost the same but the position of the peak maximum and the half-width at half-maximum (HWHM) depend on the kind of metal. The EP interphases on aluminum and copper are almost twice as wide as on magnesium and gold. In contrast to the polished metals considered above, the adhesion strength is good enough to maintain the EP–metal joint in the mechanical test. The samples fail between the metal film and the silicon wafer. After post-curing of an EP sample on evaporated copper, the sound velocity increases across nearly the whole sample to a value slightly above the maximum of the mechanical interphase prior to post-curing. A narrow interphase (ca. 80 lm) with a weak v maximum at d & 20 lm appears close to the copper (see Fig. 9.13) [12]. According to Eq. (18), the rise in v is attributed to the higher effective stiffness in the post-cured EP. On magnesium, post-curing causes similar effects in the EP layer but the details are specific for this substrate (Fig. 9.14). Inside the EP, v increases again to a value which is somewhat lower than the v maximum in the interphase after RT curing. A new interphase (ca.

9.3 Mechanical Interphases at Polymer–Substrate Interfaces Fig. 9.11 Spatial sound velocity profiles v(d) of EP (cured at ambient temperature) on Al, Mg, and Au PVD film (at d = 0 mm). The profiles of EP on Mg and Au coincide very well, whereas Al shows a different behavior.

Fig. 9.12 Spatial sound velocity profiles v(d) of EP (cured at ambient temperature) on Al, Cu, and Mg PVD film (at d = 0 mm). The profiles of EP on Al and Cu coincide very well, whereas Mg shows a different behavior.

50 lm) has formed with the maximum sound velocity at the Mg interface. All of this evidence indicates that the inhibition of the curing process (70% turnover of epoxy groups for RT curing) has been partly released within the mechanical interphase. But in the case of Mg the mechanical interphase remains even after the post-curing, and the height of the interphase has increased. For the Mg samples we investigated additionally the influence of the Si wafer on the post-cured state by removing the EP layer with the Mg from the wafer after post-curing (Fig. 9.15). This leads to an astonishing result: After release from the rigid wafer, the formerly stiffened interphase became more compliant than the bulk inside the EP layer. The width of the negative mechanical interphase exceeds 100 lm. What makes the results even more strange is that the minimum of the negative interphase is about 2% below the related interphase value of the RT-cured sample. It is obvious that DETA losses during post-curing due to diffusion are not able to explain a decrease in the sound velocity data below those values measured for the sample cured at RT.

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140

9 Mechanical Interphases in Epoxies as seen by Nondestructive High-Performance Brillouin Microscopy Fig. 9.13 Influence of post-curing on the sound velocity profile v(d) of EP on a Cu PVD film. The bulk level of the post-cured sample lies right on top of the maximum of the interphase created by curing at ambient temperature.

Fig. 9.14 Influence of post-curing on the sound velocity profile v(d) of EP on a Mg PVD film. The bulk level of the post-cured sample lies approximately on top of the maximum of the interphase created by curing at ambient temperature. Additionally the positive interphase remains and is even increased.

Fig. 9.15 Influence of the removal of the silicon wafer on the interphase behaviour of EP on Mg PVD film. The bulk level of the post-cured sample with Si wafer coincides with the bulk level after the removal of the wafer, but the formerly positive interphase turned into a negative one.

9.3 Mechanical Interphases at Polymer–Substrate Interfaces

9.3.3.3 Discussion According to the results presented in Sections 9.3.3.1 and 9.3.3.2, the adhesive interactions between the bulk metals and EP on the one hand, and between the vapor-deposited metals and EP on the other, are quite different. This can be explained by the very different surface states of the metal substrates. They differ at least in the following aspects: metal purity and morphology, surface morphology and roughness, oxidation state, and composition of the carbonaceous contamination. It is known that this difference in surface states causes drastic changes in any adhesion mechanism. For the polished bulk metals, it is concluded that the adhesion is not strong enough to change the state of the adjacent epoxy network. Another important question concerns the origin of the stiffened mechanical interphases observed in the RT-cured EP on evaporated metal substrates. These interphases extend over the very unexpected width of dozens of microns. Hence, they are not compatible with the region where the short-range chemical adhesion interactions or even physical adhesion mechanisms are at work. At first glance, a stiff mechanical interphase seems to correspond to an increased mechanical modulus. However, the epoxy shrinks during polymerization. With an adhesive bonding to the metal substrate, an internal tensile stress field builds up inevitably in the metal–epoxy interface when the epoxy vitrifies in the course of RT curing. Such a stress field may well extend over the observed interphase width and it will give a stiffening effect under the action of stress-induced polymerization. During post-curing, the epoxy is heated above the glass transition. The internal stresses will relax, at least in part, and reaction of the residual chemical groups is completed. Shrinkage due to polymerization is quite low at this stage and hence the stresses at RT will be dominated by the remaining difference in the thermal expansion coefficients of the substrate and the epoxy network. In accordance with the results for the EP layers which have been postcured on Si wafers, another stiffened interphase appears since the epoxy possesses the greater expansion coefficient. At this stage, we are unable to distinguish between the effect of stress hardening/softening and a possible increase in the interphase modulus. Fortunately, the picture is completed by the results for the samples which have been released from the Si wafer after the post-curing. Now, both the stresses from shrinking and from thermal expansion will relax quite freely because the thin metal PVD layer should be irrelevant in terms of stiffness. At RT, we expect an almost stress-free sample. Hence, the reduced stiffness in the corresponding EP/metal interphase should indicate a reduced mechanical modulus. This lower modulus would point to a particular epoxy structure in the interphase. Such a hypothesis is supported by the findings reported by Munz et al. (see Chapter 8). So the release of the inhibition of the crosslinking process seems to be caused by mechanical stress due to the bonding of the EP to a stiff substrate. We believe that these internal stress fields generated at the metal film–EP interface act in the same way as temperature does during post-curing. The stress field decays with increasing distance from the interface. If the stress field amplitude is below a certain value it is no longer able to influence the curing process and so the positive interface slows down.

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9.4 Conclusion

Using nondestructive Brillouin microscopy we have demonstrated for an epoxy system that mechanical interphases at polymer–substrate interfaces may be caused by very different mechanisms which, in turn, may yield very different morphologies. The observed interphases depend strongly on the substrate. Surprisingly, they often extend over a width of more than 100 lm. The results clearly reveal the key role of mechanical stresses in the formation of interphases. The amount of internal stress obviously depends on the adhesive interactions between substrate and polymer. Consequently, the actual surface state, including contaminations, plays an important role. Diffusion and redistribution of DETA within the DGEBA–DETA mixture, including demixing, are another important source for the formation of mechanical interphases. Having elucidated several causes of mechanical interphases, future work will have to provide more insight into the morphology and the physico-chemical properties of the mechanical interphase and their implications for the stability of the adhesive metal–polymer bonds.

Acknowledgments

This work was supported by the Deutsche Forschungsgemeinschaft. Further support was obtained from the Ministère de la Culture, de l’Enseignement Supérieur et de la Recherche du Grand-Duché de Luxembourg.

References 1 A. A. Roche, J. Bouchet, S. Bentadijne,

2

3 4

5 6 7

Int. J. Adhesion Adhesives, 22, 431–441 (2002). L. D. Landau, E. M. Lifshitz, Lehrbuch der Theoretischen Physik, Bd. VIII, Elektrodynamik der Kontinua, Akademie Verlag, Berlin, 1966. P. J. Berne, R. Pecora, Dynamic Light Scattering, John Wiley, New York, 1976. L. D. Landau, E. M. Lifshitz, Lehrbuch der Theoretischen Physik, Bd. V, Statistische Physik, Akademie Verlag, Berlin, 1966. W. Hayes, R. Loudon, Scattering of Light by Crystals, John Wiley, New York, 1978. B. Chu, Laser Light Scattering, Academic Press, New York, 1974. J. K. Krüger, in Optical Techniques to Characterize Polymer Systems (Ed.: H. Bässler), Elsevier, New York, 1989.

8 J. K. Krüger, J. Embs, J. Brierley, R. Jimi-

9

10 11

12

nez, J. Phys. D: Appl. Phys. 31, 1913 (1998). R. Sanctuary, R. Bactavatchalou, U. Müller, W. Possart, P. Alnot, J. K. Krüger, J. Phys. D: Appl. Phys. 36, 2738–2742, 2003. A. Marx, J. K. Krüger, Appl. Phys. A 47, 367, 1988. S. Dieckhoff, W. Possart, J. K. Krüger, F. Faupel, W. Brockmann et al., Adhäsionsund Alterungsmechanismen in Polymer– Metallübergängen (AAPM), Final report BMBF project, 2004. R. Bactavatchalou, Diploma thesis, Saarland University, Saarbrücken, 2003.

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10 Structure Formation in Barnacle Adhesive M. Wiegemann

Abstract

Barnacles are crustaceans that show very special adaptations to a sessile mode of life. Their adhesive is considered to belong to the most durable and toughest connections in the living aquatic world. Background information is presented on general aspects of barnacle settlement as well as on the characterization of barnacle adhesive proteins. Recent investigations of the substrate-specific supramolecular structure of barnacle adhesive and the morphology of the barnacle base are discussed. In addition to the phenomenological approach to an understanding of the adhesive properties and the structure formation processes, the concept of barnacle adhesive as a colloidal system is elucidated. The discussion of these results with observations made on other bioaquatic adherates (adhering organisms or their adhesives) reveals interesting homologies.

10.1 Introduction

Representatives from all the animal phyla living in the sea attach permanently or temporarily to solid surfaces, including those of other organisms. For some decades there has been an upsurge of interest in the adhesion mechanisms of the so-called “fouling” organisms. Adhesives secreted by aquatic organisms have to fulfill several functions, including prevention of random aggregation in the secretory glands and during transport, priming underwater surfaces, dispersion of adhesive proteins and adsorption to various materials, self-organization, and shielding from aqueous erosion and microbial degradation. Thus, studying the adhesive characteristics and the adhesion mechanisms of aquatic organisms is a promising approach toward the development of future biomimetic adhesives with an extended spectrum of properties, which may find application in the medical and dental or technical fields. Further understanding of underwater adhesion mechanisms in biological systems will also lead to conAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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10 Structure Formation in Barnacle Adhesive

clusions in the field of dehesion, which might in return promote the further evolution of nonstick surfaces employed in aquatic media (e.g., to prevent fouling of ship hulls or medical implants). Barnacles are crustaceans that show very special adaptations to a sessile mode of life. Acorn barnacles that are cemented to the substratum with the whole calcareous or noncalcareous base are frequently found on ship hulls – most of them belong to the family Balanidae. The adhesive of these barnacles is considered to be among the most durable and toughest connections in the living aquatic world [1], and it is therefore also called “cement”. The following synopsis summarizes the results of the investigations by Wiegemann and Watermann [2] of the subsrate-specific supramolecular structure of barnacle adhesive and the morphology of the barnacle base. In addition to the phenomenological approach to understanding the adhesive properties and the structure formation processes, the concept of barnacle adhesive as a colloidal system is elucidated. The discussion of these results with observations made on other bioaquatic adherates (adhering organisms or their adhesives) reveals interesting homologies.

10.2 Barnacles 10.2.1 General Aspects of Barnacle Settlement

Barnacles produce pelagic larvae (nauplii) that undergo six developmental stages before they reach the settling stage (cypris larvae). The cyprid’s search for a suitable attachment site is stimulated by a multifactorial spectrum in which physical and chemical surface characteristics play a part [3, 4]. Permanent settlement is initiated by the secretion of a small amount of adhesive in which the attachment organs of the cyprid’s antennules are embedded. Thereafter the cyprid undergoes metamorphosis to a juvenile barnacle. It may take several days until the first post-metamorphic adhesive is secreted. Usually, this adult cement appears as discrete rings underneath the relatively even base plate of the barnacle. The secretion is most likely linked to the molt cycle since the cement-producing cells are modified epidermal cells [5]. During the growth process new adhesive is released through pores close to the margin of the shell. The liquid matrix fills the gaps between the base plate and the surface by capillary action and excludes water at the same time. In the case of Balanus eburneus the adhesive cures within 6 h [6] to a rigid mass (hereafter called cement).

10.2 Barnacles

10.2.2 Biochemical Characterization of Barnacle Cement

The adhesive is a proteinaceous material (> 90% protein), while the remainder consists of carbohydrate, ash, and trace amounts of lipid [5]. Based on the assumption that the cement glands are homologous with the glands that secrete the cuticle [7] (the cuticle of arthropoda is tanned by o-quinone crosslinking) and due to the existence of phenols and polyphenol oxidase in the cement glands of the settling stage it was widely postulated that quinone crosslinking also occurs in the cement [5]. The comprehensive work on the cement glands of the settling stage of the barnacle clearly concluded that cypris cement contains tanned protein, and hence quinones [8]. Despite extensive histological and histochemical studies on the cement glands of adult barnacles (see Refs. [9–11]) a similar picture did not emerge. The extreme resistance of the hardened cement toward salt solutions, dilute acid, and alkali suggested that hydrogen or salt bonds were too weak to be the main crosslinking mechanisms. Disulfide bonds were also precluded due to the resistance of the adhesive to thioglycolate (see Ref. [12]). In contrast, Barnes and Blackstock [13] and Yan and Pan [14] showed that an anionic detergent, sodium dodecyl sulfate (SDS) containing the reductant 2-mercaptoethanol (2-ME) is sufficient to dissolve the cement, what lead to the assumption that hydrophobic interactions and sulfur crosslinks are key components of the cement matrix. Though Naldrett and Kaplan [15] found high amounts of the amino acid tyrosine (Tyr) in a major protein of barnacle cement, which is modified to DOPA in the mussel adhesive, infrared and NMR spectroscopy provided evidence that quinones are absent from the cement of balanid barnacles [12, 15]. Still, other post-translational modifications were reported to be present, possibly being glycosylations [15]. It was suggested that the highly hydrophobic character of the residues was the main reason for the insoluble nature of these proteins. In more recent studies it was possible to render soluble at least 90% of the cement of the barnacle species Megabalanus rosa [16, 17] and Balanus eburneus [15]. The authors found a complex of distinct proteins, which achieve the adhesion cooperatively. It was reported that these proteins are crosslinked through cysteine residues. In contrast, a new study [18] showed that the presence of a reducing agent had only a minor impact on the solubility. The strong effect of the denaturing conditions (the presence of sodium dodecylsulfate (SDS), elevated temperature) on the solubility of the barnacle cement suggests a complex structure of the barnacle cement noncovalently bonded. The suggestion was reinforced by Raman spectra, which did not indicate the presence of these groups in cement of the species Balanus crenatus. The authors [18] additionally found a triplet of proteins, which further broke down into subunits under denaturing conditions. It was suggested that this protein complex in its quaternary structure formed by small subunits is equivalent to the globular cement structures detected previously [2] (see below).

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10.2.3 Substrate-Specific Formation of Barnacle Adhesive

It is very well known that the type into which a cell is differentiated is determined by the interaction of the cell surface with its (macro)molecular environment (e.g., Refs. [19, 20]). It is shown below that the supramolecular structure of barnacle adhesive, which is an extracellular matrix, is also formed through surface interactions. Investigations by ex-situ scanning electron microscopy (SEM) and in-situ atomic force microscopy (AFM) have recently shown that barnacle cement is composed of globular structures of nanometer size [2]. These formations might be comparable with the granular structures detected in AFM studies by Berglin and Gatenholm [21] observed on a poly(dimethylsiloxane) (PDMS) surface. Wiegemann and Watermann [2] further reported that these granules formed specific supramolecular structures in response to certain substrate characteristics: The cement of well-adhering barnacles was very dense, appearing as a thin sheet or a sponge-like matrix (e.g., adhesive secreted onto aluminum, Fig. 10.1 a). The

Fig. 10.1 SEM images of various appearances of the supramolecular structure of barnacle adhesive: (a) thin sheet of densely arranged nanometer-sized adhesive globules on aluminum foil; (b) sponge-like appearance of adhesive with the tendency to form strands on a self-polishing coating; (c) weblike adhesive of loosely matted strands on a PDMS coating.

10.2 Barnacles

tendency to form adhesive strands (Fig. 10.1 b) on substrata (e.g., a self-polishing coating) on which barnacles achieved lower adhesion strength, resulted in the creation of branched or loose web-like “cement” structures (Fig. 10.1 c) on nonstick coatings based on PDMS. Compared with the structures described above, the latter were highly hydrated and of reduced cohesive strength. 10.2.4 Substrate-Specific Morphology of Barnacle Base

Barnacles usually produce an even base plate that nestles against the substrate surface. However, barnacles that were loosely attached to PDMS coatings produced cup- or funnel-shaped base plates filled with a hydrated adhesive plaque. Cross-sections through the barnacle base viewed with a light microscope revealed that the adhesive plaque was a multilayered system, in which the density of the adhesive layers decreased consecutively toward the substrate (Fig. 10.2). Cup-shaped base plates were not observed only on PDMS coatings; barnacles are also able to grow into soft material that lacks nonstick characteristics, thus producing cup-shaped base plates [22]. Generally, the base plate enlarges horizontally while the parietal plates grow downward (Fig. 10.3). Materials with relatively low cohesive strength (e.g., wood, or soft conventional coatings) are cracked by the growing shell. With each growth increment the shell of the barnacle sinks deeper into the soft substratum. However, the barnacles were not able to crack the PDMS coating. It was assumed that the adhesion strength was too weak to hold the barnacles in place while the parietal plates exerted downward pressure onto the substratum [2]. Consequently, the loosely attached barnacles were lifted, and due to the horizontal and downward apposition of lime the base plates grew into cup shapes. If a barnacle lifted itself, the freshly released adhesive would consequently flow into the emerging space. A series of such events would cause the adhesive to be built up layer by layer (see the scheme in Fig. 10.4 for the deformation of the barnacle base and the location of adhesive secretion).

Fig. 10.2 Cross-section of a young barnacle that has been removed from the PDMS coating (hematoxylin- and eosin-stained). The space underneath the cup-shaped base plate (B) is filled with cement (C), which is arranged in layers. Within the adhesive plaque the staining is most intense close to the base plate (o C indicates the oldest cement layer), fading with each consecutive layer.

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Fig. 10.3 Schematic cross-section through a barnacle showing the growth directions of the barnacle base according to Gutmann [10]: I: downward growth of the parietal shell plates; II: horizontal enlargement of base plate.

Fig. 10.4 Schematic cross-section through a barnacle grown on a PDMS coating with a cup-shaped base plate and multilayered adhesive plaque. Cement ducts open at the margin of the barnacle base, from where the cement will flow into the interfacial gaps between the substrate and the base plate.

10.2.5 Phenomenological Approach to Adhesive Structure Formation and Morphology Changes

Thick gel-like adhesive was detected whenever there was a large distance between the barnacle base and the substratum to be bridged (e.g., on fiber coatings or on surfaces with thick biofilm layers). It is questionable whether the barnacle produces more adhesive in order to continuously fill out the space to the substratum or whether the adhesive is expanded by water uptake. As the cement-producing cells appeared normal, water uptake is more likely to be the dominant process.

10.2 Barnacles

It was proposed that water uptake is a general mechanism in barnacle adhesive for the purpose of expansion [2]. This mechanism occurs whenever large distances between the barnacle base and the substratum are to be bridged and if an adequate amount of adhesive is not available. It is possible that the adhesive proteins that need to be distributed all over the gap cannot form a tight net because of the large distances that have to be bridged. Instead the gaps are filled by expansion of the adhesive through water uptake. As a barnacle attached to a PDMS coating grows, the gap produced underneath the cup-shaped base plate increases. The increasing gap offers more space for the adhesive to expand and to take up water, which results in a continuous reduction of density and cohesive strength within the multilayered adhesive plaque, as described above. There is an apparent difference in the ability of barnacle adhesive to take up water when secreted into small gaps or into large gaps. The key to this contradiction might lie in the capacity of proteins to macroconformational changes organizing their hydrophobic and hydrophilic groups into regions, hiding the former or the latter in the interior according to the properties of the aqueous environment (see also Section 10.4). Though the swelling process reduces the number of potential crosslinking interactions causing low cohesion, the mechanism provides benefits: · Gaps between base plate and substratum are filled without an extraordinarily high production of adhesive, which would be a burden to the barnacle’s energy resources. · The hydrated adhesive is more flexible than strongly crosslinked adhesive. Flexibility is a useful mechanical property for an adhesive to bond to a mobile or elastic substratum (e.g., PDMS, biofilm). The web-like structure of the hydrated adhesive is rather more able to adjust to movements of the substratum (molecular mobility of the polymer backbone, compression or extension of elastic material) than the rigid, strongly crosslinked cement. · Another advantage of the reduction of the cohesive strength of the adhesive might also be seen in the persistence of the tackiness, which could be similar to the principle realized in pressure-sensitive adhesives, e.g., tapes [23]. From the evolutionary point of view, structures or mechanisms that prove a success are likely to be passed on. However, adult barnacles probably do not control the structure formation process of their cement. The ability of cypris larvae to sense physical and chemical surface properties most likely vanishes after attachment and metamorphosis, for the main function of the adult is to feed and to metabolize the energy into growth and reproduction. Instead, the above observations probably suggest that the structure formation process is simply a result of water uptake increasing the effective size of the adhesive filling spaces or gaps.

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10.3 Homologous (?) Structure Formation of Biological Adherates on Hydrophobic Surfaces

Adhesion is essential to the living world and is commonly found in any creature. Thus, it should not be surprising to find corresponding or even homologous mechanisms fundamental to adhesion. Bioadhesives may basically consist of proteins, polysaccharides, polyphenols, and lipids. These substances occur mostly in combination. The adhesives of the well-studied blue mussel (Mytilus edulis) and of barnacles are proteinaceous materials. Other well-known proteins with adhesive properties are elastin, collagen, fibronectin, laminin, fibrinogen, and keratin. In any case, the quality of adhesion is very much linked to the dispersion of the liquid phase and the wettability of the substrate surface. Liquids will generally spread freely on solids of high surface free energy (also termed hydrophilic) but might show nonspreading behavior on solids of low surface free energy (also termed hydrophobic). Thus, it could be expected that organisms adhere most strongly to high-energy surfaces. It has been confirmed in numerous studies that organisms adhere more strongly to hydrophilic surfaces than to hydrophobic surfaces (e.g., Refs. [24–26]; specifically for M. edulis [27]; specifically for barnacles, e.g., Refs. [28, 29]). Other experiments revealed a depression of fouling within the low surface tension range of 20–25 mN m–1 (e.g., Ref. [30]). This phenomenon is also known as the “biocompatible range” or “Baier’s Window”, for which Baier [31] and Dexter [32] gave a thermodynamic explanation: The surface energy within the biocompatible range has a minimum interfacial tension with aqueous solutions of biological macromolecules, and gives the minimum driving force for adsorption. (In order to prevent confusion: although such surfaces repel adhesive biomolecules and are hence incompatible, these surfaces are biocompatible when used as implant material for medical purposes.) Yet according to the literature there are some larvae which preferentially settle on hydrophobic substrata, e.g., the bryozoan Bugula neritina Linnaeus [33] and the ascidian Ascidia nigra [34]. It is known that such organisms colonize substrata successfully despite lower adhesion [35]. My own observations regarding bryozoans suggest that these organisms may profit by a selective advantage (e.g., a favorable hydrodynamic shape) in spatial competition for low-energy surfaces, which enables them to occupy a rather extreme niche. The adhesion strength of bryozoans, however, seems to be higher on hydrophilic surfaces (e.g., on barnacle shells). In addition to the low adhesion strength of organisms on hydrophobic surfaces, it is even more intriguing that secretions on such surfaces showed an increase in volume in numerous cases (see also Table 10.1): Investigations of M. edulis revealed that the area or volume of the adhesive plaque is influenced by the surface energy of the substratum: on more polar surfaces (slate and glass), smaller plaques are formed [25]. Similar results were reported specifically for

10.3 Homologous (?) Structure Formation of Biological Adherates on Hydrophobic Surfaces Table 10.1 Some peculiarities of adherates on hydrophobic surfaces.

Adherate

Observation on a hydrophobic surface in comparison with a hydrophilic one

References

Barnacle adhesive

increas e in volume (water uptake), low adhesion strength

[2, 21, 29]

increase in volume, low adhesion strength increase in volume increase in volume, low adhesion strength increase in volume increase in volume

[25] [27] [36] [37]

M. edulis Adhesive plaque Adhesive protein Mefp-1 Conditioning films Human salivary films Algae rhizomes Fucus Enteromorpha Ulva Various macroalgae Bacteria colonies

morphological changes

[38] [39] [24, 40] [41] [42] [43]

the M. edulis foot protein (Mefp)-1, which produced a hydrogel-like film on a hydrophobic surface whereas on a hydrophilic SiO2 surface a compact rigidly attached protein layer was formed [36]. Even conditioning films, which are macromolecules associating quickly on a surface after submersion in an aqueous medium, were reported to be flatter and tightly adhering to surfaces with high surface free energy, while on surfaces of low surface free energy they were weakly bound but thicker [37]. Similar observations were made by Vassilakos et al. [38], who observed that salivary films absorbed on hydrophilic surfaces were thinner than those on more hydrophobic ones. Morphological differences depending on the surface tension were also described for cell complexes, in particular for the rhizomes of algae (for Fucus [39], for Enteromorpha [24, 40], for Ulva [41], and for various macroalgae [42]). The algae developed a compact discoid rhizoidal base with short, tight-jointed filaments on hydrophilic surfaces, while on hydrophobic surfaces the filaments were long and free, giving up to a threefold increase in the algal base area. Dalton et al. [43] even reported that the substrate surface hydrophobicity influences the morphology of bacteria colonies. The observations reported above may raise the question of whether there is a common mechanism inducing a swelling process of extracellular matrix adhesives or cellular adherates on hydrophobic surfaces. It can be speculated that the poor wettability of hydrophobic surfaces is a result of repulsion. Fant et al. [36] specifically suggested for M. edulis adhesive that charged polar groups – making up the majority in Mefp-1 – are probably highly hydrated in aqueous environments and therefore repelled from the surface. It can further be suggested that the repulsion would hamper particles or

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molecules of potential adherates from depositing onto the surface and from aggregating, and it would cause a diffuse distribution of the adhering matrix leading to a loose formation with high water content.

10.4 Theoretical Colloid Approach to Structure Formation in Barnacle Adhesive

As reported above, the stage of expansion through water uptake in barnacle adhesive is associated with specific supramolecular structures of the adhesive matrix. Remarkably, foam-like adhesive and globular structures such as those observed for barnacles [2] were also found in mussel adhesive [44]. The authors suspected the globules observed in the adhesive plaque of M. edulis to be caused by incompletely dispersed matrix concentrated in the form of densely packed globules. The assumption partially agrees with the suggestions of Wiegemann and Watermann [2] concerning their observations in barnacle adhesive. The distribution of barnacle adhesive proteins and therefore the density of the adhesive were observed to vary greatly within the cement plaque and also to depend on the substrate. Generally, dense adhesive matrix that coincided with high cohesive or adhesive strength appeared as a thin sheet at the interface to the environment (substrate or aquatic medium). The M. edulis adhesive plaque also possesses a highly crosslinked superficial layer, which was attributed to the cuticular protein Mefp-1 [45], protecting the vulnerable matrix proteins Mefp-2 and Mefp-4 [45, 46], which form a solid foam in the interior of the plaque [44]. In contrast, the decrease in density in barnacle adhesive, which was manifested in foam-like structures or – in the extreme – in loosely matted strands, was suspected to be caused by expansion through water uptake [2]. The process of structure formation of barnacle adhesive can be further discussed in the light of a biocolloidal system: The characteristics of biocolloids are based on the principle of self-organization [47]. Proteins are macromolecules which carry a mixture of polar and nonpolar side chains, and they are able to organize their hydrophobic and hydrophilic groups into regions, hiding the former or the latter in the interior according to the properties of the aqueous environment. Thus, barnacle adhesive proteins could be considered as surfactants, which are known to aggregate in dispersion into micelles or at interfaces into ordered surfactant films. This property causes the migration of surfactants from the bulk of the dispersion to the interface, where they accumulate or are adsorbed to a foreign surface. The high concentration of proteins in the periphery of the adhesive plaque of barnacles or mussels might be caused by a similar process. Single molecules or micelles associate spontaneously in a thermodynamic equilibrium at a definite critical micelle concentration within a biocolloidal system [47]. Analogously to micelle formation in liquid systems, aggregation of surfactants at a surface depends on a critical hemi-micellar concentration [48, 49]. The removal of the hydrophobic molecular region from the hydrophilic interface

10.4 Theoretical Colloid Approach to Structure Formation in Barnacle Adhesive

(or, vice versa, the removal of the hydrophilic region from the hydrophobic surface) is an energetic driving force of this process. It may be conceivable that the structure formation of barnacle adhesive is determined by critical self-assembly concentrations of the adhesive proteins within an interfacial gap between a barnacle base and a substrate. It can further be suggested that the biopolymers form coherent gel structures, in which two transitions of critical protein concentrations determine the arrangement of adhesive globules: from a dense sheet-like formation to a slightly loose sponge-like formation to a very loose branched or web-like structure. The consideration outlined above agrees with the assumption that substrates influence the structure formation of barnacle adhesive indirectly via properties that hold the barnacle base at a distance from the surface (e.g., fiber coatings, thick biofilm layers) or cause large interfacial gaps (e.g., PDMS coatings). These factors affect the concentration in the dispersion, and hence the “position on the adsorption isotherm” as compared to the micellar or hemi-micellar concentration, by impinging on the interfacial volume that needs to be filled with adhesive. By relating these thoughts to hydrophobic surfaces inducing a swelling process of adherates, the following model can be proposed: The repulsion causes a diffuse distribution of the adhering matrix, which may spread widely (e.g., into the aquatic medium in which the organism lives, or within the large interfacial gap between the cup-shaped base plate of a barnacle and the substrate surface). Consequently, the repulsion would cause a low concentration of adhesive matrix within the bulk phase, which would lead to formation of a loose structure with high water content. An important question has occurred often in the literature regarding the trigger for the curing of barnacle adhesive. It was speculated until now that the adhesive is a two-component system [50], the components being separated by enclosure into micelles. The micelles and their contents are thought to fuse after secretion, which initiates polymerization. In continuation of the theoretical colloid approach above, it appears possible that protecting colloids or emulsifying agents prevent the adhesive proteins from aggregation or fusion similarly to the situation in adhesive dispersions. Since the molecules do not interact at this stage, their molecular weight does not influence the viscosity (in contrast to solvent-based adhesives, in which the molecules interfere with each other with increasing concentration [23]). This would also explain the low viscosity of freshly secreted barnacle adhesive as reported by Walker [5]. The thermodynamically pseudo-stable state, which is maintained through the action of protecting colloids or emulsifiers, could be disrupted by modifications of conditions within the dispersion (e.g., the pH) or by mechanical impacts. Such changes could trigger the gel formation. A promising approach toward the unraveling of bioaquatic adhesion mechanisms would therefore be to search for such emulsifiers and the identification of the triggers inducing the curing process, while continuing the characterization of the adhesive matrix. A long history of references to the characterization of the adhesive protein complex of barnacles (see Section

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10.2.2) and especially the adhesive proteins of M. edulis (e.g., Refs. [51–59]) already exists in the scientific literature (for a current review of the characterization of barnacle and mussel adhesive see Ref. [60]).

10.5 Conclusions

Knowledge of the structure of barnacle cement within colloidal dimensions has recently been gained. It is shown that the barnacle adhesive presents a useful tool for understanding and investigating formation of bioadhesive structures as colloidal systems. From the observations regarding the swollen condition of adhesives or adherates absorbed to hydrophobic surfaces, it can be concluded that the relevant basic processes might refer to general mechanisms in the living aquatic world. Even the aggregation of globular adhesion proteins to specific supramolecular structures observed in barnacle cement may be speculated to be a general theme in aquatic organisms, since similar structures have been found in mussel adhesive.

Acknowledgments

This work was kindly supported by the Bonding Technology and Surfaces Department of the Fraunhofer Institute for Manufacturing and Advanced Materials in Bremen. Special thanks go to Dr. M. Noeske for fruitful discussions, inspirations, and comments on the manuscript.

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11 Adhesion Molecule-Modified Cardiovascular Prostheses: Characterization of Cellular Adhesion in a Cell Culture Model and by Cellular Force Spectroscopy U. Bakowsky, C. Ehrhardt, C. Loehbach, P. Li, C. Kneuer, D. Jahn, D. Hoekstra, and C.-M. Lehr

Abstract

In vascular surgery the insertion of artificial blood vessels is a common method, but limited to specific medical applications. Long-term studies have revealed that unspecific adhesion of fibrin and collagen to the implant surface leads to thrombus formation and occlusion, even though PTFE is extremely inert and hydrophobic. In the present study we describe the modification of cardiovascular prostheses made of PTFE by a chemical method in combination with a biochemical method, leading to improved cell adhesion. The surface was covalently coated with extracellular matrix-adhesion molecules and with a fragment of the bacterial invasin protein of Yersinia enterocolitica, the so-called Min3. All these molecules showed high binding constants for a4–b1-integrins, which were expressed on the surface of endothelial or fibroblast cells. On this protein-modified surface, primary endothelial cells (HUVECs) were seeded in the presence of growth factors, and cells were exposed to physiological shear stress after an initial adhesion period. To characterize the adhesion properties of the HUVECs, we used scanning force microscopy (SFM). In this study, we developed a method to determine the specific interactions of HUVECs with various adhesion molecules using SFM. HUVECs were grown at the edge of a “tipless” cantilever. We observed higher cell adhesion rates and higher adhesion forces when using Min3 than for collagen or fibronectin. The measured adhesion force depended on the cell surface area and the contact time (between the implant and the cantilever), and ranged between 0.2 and 1.2 nN.


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11.1 Introduction

Cardiovascular disease is the primary cause of death, in both the United States and Europe, with arteriosclerosis being the most common form. The process of arteriosclerosis is considered to consist largely of the accumulation of lipids within the artery wall [1]. By surgical therapy new pathways (conduits) are being constructed, connecting the aorta or other major arteries and distal segments of arteries (e.g., coronary arteries) with vessels beyond the stenosis or obstructing lesions. By this means, blood supply to ischemic tissues can be re-established [2]. Reversed segments of autologous saphenous vein can be used as appropriate conduits. Particularly in coronary revascularization, the internal mammary arteries and the radial arteries have widely been used [3]. However, autologous vessels may be insufficient for multiple or repeated bypasses and/or saphenous veins may have varicose degenerative alterations. Therefore, allograft arteries and veins as well as synthetic tubes have been developed, but they proved to be less satisfactory as conduits [4, 5]; particularly, synthetic grafts with an internal diameter of 6 mm or less are prone to thrombus induction and occlusion [6]. In this context, the development of novel biocompatible materials becomes more and more important. Current grafts are restricted to large and medium diameters since none of the available ones is suited for the replacement of smalldiameter vessels or coronary arteries, respectively. A slight blood flow favors the formation of thrombi, resulting in a subsequent occlusion of the vessel. By administration of anticoagulative drugs, this outcome can be reduced, but over the years the formation of a neointima occurs, which consists of both stromal cells and endothelial cells. In particular, stromal cells are prone to hyperplasia, thus steadily reducing the vessel lumen, which results eventually in the formation of a stenosis or the occlusion of the prosthetic graft. Local clotting seems to be the central mechanism of thrombus formation in arteriosclerosis of native blood vessels and limits the performance of prosthetic grafts [7–11]. The principle of in-vitro endothelialization of vascular grafts has been established and clinical results are available [12–17]. In this case, a lining of vascular grafts with endothelial cells is possible, but a confluence exceeding 70% of the inner surface, can only be reached by the use of growth factors in the culture medium. Apparently, this limitation in cell density is the reason for the fact that present applications with endotheliazed grafts show no significant refinement [14, 18–23]. One explanation for these difficulties is the small cell yield obtained from isolation of donor vessels; it can be increased by addition of growth factors, but then cells might undergo dedifferentiation in batch culture. Another problem occurring with dedifferentiated cells is the adherence of leukocytes to their surface, which in turn leads to the formation of a thrombosis, resulting – in the worst case – in an occlusion of the vessel. In addition to that, insufficient adherence of endothelial cells to the vascular graft, caused by the inertness and hydrophobicity of the polymer wall material, leads to more than 80% of the cells lining the graft being washed away when physiological shear stress is applied. Two biomaterials dominate the vascular

11.1 Introduction

graft market to date, Dacron (PET, poly(ethylene terephthalate)) and PTFE (polytetrafluoroethylene). Dacron monofilaments are woven or knitted into various designs to form the graft. They are generally used in the large-diameter category (12– 22 mm). Expanded PTFE consists of nodes and fibers and was generally used in the intermediate-diameter category (6–12 mm). The porosity of these PTFE grafts encourages the formation of a biological lining on the luminal surface, known as neointima (cellular lining) or pseudointima (acellular lining). Therefore, modified PTFE grafts are more suitable candidates for small-vascular prostheses than PET. The concept of endothelial cell seeding on vascular implant surfaces was developed to mimic physiological blood vessels and to further reduce unspecific adhesion [11, 24]. PTFE, however, is a very hydrophobic matrix and thus not the optimum substrate cell culture. Adhesive coatings with aqueous proteins on the implant surface have given only inadequate performance under shear stress conditions [25–27]. Therefore, covalent attachment of specific adhesion molecules may provide a means to achieve a strong, shear stress resistant, cell binding. In this study, we describe the modification of PTFE by a chemical method in combination with a biochemical method leading to an improved cell adhesion. The modification was first developed with punched disks from PTFE film, and was then transferred to commercially available vascular grafts made of PTFE. The aim was to retain the mechanical properties of the polymer so that only the biological properties of the inner surface of the graft were changed. Thus it would become possible to couple adhesion proteins to the previously inert and hydrophobic polymer surface to favor adhesion and attachment of endothelial cells (ECs) withstanding shear stress as it occurs in small-diameter vascular grafts. We tested a three-step-preparation in order to alter the polymer surface selectively. In the first of these steps, reactive groups should be formed on the inner surface of the hydrophobic PTFE, which than make it possible to couple a crosslinker for further attachment of the adhesion molecules. Reactive surface groups such as hydroxyl, amino, or carbonate can be generated on formerly inert PTFE surfaces by different methods, including dry-chemical treatment (e.g., plasma etching, ablation) or wet-chemical methods (e.g., treatment with H2O2/H2SO4, chromate, permanganate) [28–31]. The activation of PTFE tubes with H2O2/ H2SO4 as oxidant was our first choice, because it resulted in a sufficient number of free functional groups without damaging the structural integrity of the polymer combined with good handling. In a second step, the crosslinker, cyanuric acid, was covalently bound via hydroxyl groups to the activated PTFE. Cyanuric acid theoretically provides two binding sites, hydroxyl and amino groups, for proteins. The family of extracellular matrix (ECM)-proteins (e.g., laminin, collagen, fibronectin) and their derivates (RGD peptide sequences) came into consideration for the final functionalization of the activated and crosslinked polymer. Here, we struck out on a new path by choosing a fragment of the bacterial invasin protein of Yersinia enterocolitica. This bacterium enters mammalian organisms via the gut and the protein that enables this access is called invasin A. Primary human umbilical venous epithelial cells (HUVECs) were seeded on this protein-modified surface. Culture was performed in the presence of growth

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factors and cells were exposed to physiological shear stress after an initial adhesion period. To characterize the adhesion properties of the HUVEC, we used scanning force microscopy (SFM). In this study, we developed a method to determine the specific interactions of human endothelial cells (ECs) with various adhesion molecules using SFM. The ECs were grown at the edge of a tipless cantilever. The adhesion molecules (i.e., gelatin, fibronectin, collagen, Min3) were covalently coupled to an artificial blood vessel made from PTFE. Due to unspecific adhesion of fibrin and collagen to the implant surfaces, thrombus formation and occlusion may occur in clinical use. To overcome this problem, the concept of EC cultivation on modified implant material was developed to mimic the physiological surface of blood vessels [36]. The expiration of these cell layers strictly depends on the cell-surface interaction, which can be improved by modifying the graft surface with adhesion molecules.

11.2 Materials and Methods 11.2.1 Chemicals for the Modification

H2SO4 and chloroform were obtained from Merck, H2O2 from Roth, and cyanuric acid from Sigma. All chemicals used for the modification were of analytical grade. Collagen, fibronectin, gelatin, and BSA (bovine serum albumin) were purchased from Becton Dickinson. The adhesion molecule from Yersinia entericolitica, Min3, was expressed in Escherichia coli, isolated, and purified by D. Jahn’s group (Universität Braunschweig). 11.2.2 Implant Materials

Commercially available PTFE was used for the investigations (Merck). We developed the method and the parameters (acids, concentration, ratio, time) for the wet-chemical modification of PTFE on the basis of PTFE film material (thickness 0.1 mm) punched in 12 mm disks. Then the results from the PTFE film model were transferred and adjusted to medical grade vascular grafts (B. Braun AG; see Fig. 11.1). 11.2.3 Modification of the PTFE Surface

In a first reaction step, the surface was treated with H2O2/H2SO4 (1 : 1) at room temperature for 20 min. In the second step, the reactive hydroxyl groups generated were esterified with cyanuric chloride with release of HCl. For activation,

11.2 Materials and Methods Fig. 11.1 Cardiovascular prosthesis used for surgical treatment of an aneurism of the aorta in a real patient. Courtesy of Prof. Dr. H.-J. Schaefers (Homburg/Saar, Germany).

we used a filtered 10–3 m solution of cyanuric chloride in chloroform for 1 h followed by intensive washing with chloroform to remove excessive cyanuric chloride (Figs. 11.2, 11.3). In the third step, the activated PTFE surface was finally coupled with functional molecules to achieve targeted adhesion and spreading of endothelial cells. High affinity and specific interactions of these molecules with physiological receptor molecules on the cells were of importance. Candidate adhesion molecules are constituents of the extracellular membrane (collagen, fibronectin, lami-

Fig. 11.2 Reaction scheme of the wet-chemical modification of PTFE. (A) unmodified PTFE; (B) oxidized PTFE with possible reactive groups formed; (C) binding of the crosslinker cyanuric chloride; (D) covalent binding of adhesion protein to PTFE.

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11 Adhesion Molecule-Modified Cardiovascular Prostheses Fig. 11.3 Schematic of PTFE modified with Min3.

nin, vitronectin), other peptides and proteins containing the RGD motif, or adhesion proteins of microorganisms such as the invasin A of Yersinia species, and derivational sequences. These adhesion proteins target specific receptors on the cells, especially b1 and b3 integrins. In earlier experiments, we observed a higher cell adhesion rate during 24 h of incubation when using the bacterial invasin (Min3) than with collagen or fibronectin. This molecule binds specifically and with a high affinity to a4–b1 integrins (vla-4) [15]. The association constant is approximately three magnitudes higher than for the commonly used extracellular matrix proteins (collagen, fibronectin) and RGD-peptides. Min3 was produced by heterologous expression in E. coli and purified by maltose-binding affinity chromatography and ion-exchange chromatography to a purity of > 95%. After intensive washing with chloroform, an aqueous solution of the test adhesion protein (borate-buffer, pH 8.4) was incubated for 12 h to allow reaction of surface-bound cyanuric chloride residues with amino groups of the protein. 11.2.4 Scanning Force Microscopy

SFM was performed on a Nanoscope IV Bioscope (Veeco). The microscope was vibration- and acoustic-damped (Brunhild II, Veeco). Commercially available pyramidal Si3N4 tips (Veeco) on a V-cantilever (length 125 lm) were used for the force measurements. For the force measurements with attached ECs, the tip was removed mechanically from the cantilever. The measurements were done in contact mode with PBS (pH 7.4) medium. The force constants of the cantilevers were determined according to Florin et al. [37] and were in the range 0.06–0.14 N m–1. The scan speed was proportional to the scan size and the scan frequency was between 0.5 and 1.5 Hz. Images were obtained by displaying the height signal and friction force signal, both signals being recorded simultaneously. The interaction between the cantilever-attached ECs and the proteinmodified PTFE surface was determined by recording and analysis of the force– distance curves. For the attachment of ECs to the nanosensor, the cantilevers were pre-activated with cyanuric chloride and modified with Min3. The cells were cultured on the cantilever for three days post-seeding, used for the experiment, and afterwards analyzed by scanning electron microscopy (SEM). Cells were successfully attached on three out of ten cantilevers.


The visualization of the protein modification was performed under normal atmospheric conditions, at a temperature of 25 8C and with a humidity of approx. 60%. We used commercially available pyramidal Si3N4 tips (NCH-W, Digital Instruments, Santa Barbara, CA) in tapping mode with a scan frequency between 0.5 and 1.5 Hz, a resonance frequency of 220 kHz, and a nominal force constant of 36 N m–1. 11.2.5 Fourier Transform Infrared Spectroscopy

The surface modification was further followed by FTIR (ATR) spectroscopy on a Perkin-Elmer System 2000. Analysis was performed with an internal reflection element (GaAs crystal, 458 incidence angle) with 100 scans co-addition and a resolution of 4 cm–1. The insoluble, inert PTFE was cut into thin microtome slices between 150–500 nm and analyzed in transmission mode. 11.2.6 Environmental Scanning Electron Microscopy

The ESEM (XL30 ESEM-FEG, Philips) that was used warranted a lateral resolution of 3.6 nm and was therefore qualified to visualize the global morphological appearance of the polymer surface and to characterize the cell-modified nanosensors (cantilevers). The unmodified and chemically activated specimens were sputtered with platinum without any other treatments. The cell-covered specimen (cantilevers and PTFE) were fixed with 2.5% glutaraldehyde–phosphate-buffered saline solution (Sörensen/Arnold; pH 7.4) for 2 h and, after extensive washing, drained in an ascending ethanol column (30, 50, 70, 80, 90, 96, and 100%). After this procedure, the specimens were transported in water-free acetone to remove alcohol, subjected to critical point drying, and sputtered with platinum. All specimens were investigated under vacuum conditions (10–3–10–4 atm). 11.2.7 Confocal Laser Scanning Microscopy (CLSM)

For CLSM, PTFE grafts lined with endothelial cells were treated as follows. Specimens were fixed in 4% paraformaldehyde (PBS, pH 7.4) for 15 min, washed with PBS, and permeabilized with a 2% Triton X-100 solution in PBS for another 15 min. After washing again with PBS, they were incubated with the primary monoclonal antibody vWF (CD31) (DAKO) in a 1 : 50 dilution at 37 8C for 30 min. After another washing step, the specimens were incubated with the secondary antibody (FITC-labeled, DAKO) in a 1 : 100 dilution at 37 8C for 30 min. Then, cell nuclei were stained with a 1 : 1000 solution of propidium iodide at 37 8C for 15 min. After extensive washing and a short air drying, the PTFE matrices were mounted onto glass slides with FluoroSave (Calbiochem)

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and hardened overnight. The specimens were studied with an MRC1024 confocal laser scanning microscope (BioRad). 11.2.8 Isolation and Culture of HUVECs

HUVECs were isolated within 6 h of childbirth [38]. During transport and processing, umbilical cords were stored in sterile HBSS (Hanks’ balanced salt solution) supplemented with an antibiotic–antimycotic solution (Sigma). Ensuring sterility, 1 cm lengths of both ends of the umbilical cord were cut off and the cord was cannulated immediately from both ends with a nub canula (Medica) and kept in place with mosquito clamps (Medica). Subsequently, the vein was perfused with sterile HBSS until the effluent was free of red blood cells. After the vein had been depleted of air, it was filled with 0.25% collagenase type I solution (Worthington) prewarmed to 37 8C and plugged with Luer lock plugs. The cord was then incubated for 15 min at 37 8C. For increased cell yield, the cord was gently massaged and the vein was flushed with 50–100 mL of 10% FCS (Foetal Calf Serum) in HBSS. The collected effluent was centrifuged at 500 g for 10 min. The cell pellet was washed twice with HBSS and centrifuged again at 500 g for 10 min. The cells were resuspended in EGM II culture medium (Promocell), filtered through a piece of sterile gauze, and seeded into 75 cm2 culture flasks (Greiner) freshly coated for 30 min with 0.1% gelatin in phosphate-buffered saline (PBS). Cell cultures were incubated at 37 8C in a 5% CO2 atmosphere. The medium was changed every second day. When they were approaching 70–80% confluence, cells were harvested using trypsin/ethylenediaminetetraacetate (EDTA) (Sigma) and passed onto gelatincoated culture flasks at a split ratio of 1 : 4. 11.2.9 Endothelialization of PTFE Films

Cell seeding experiments were performed after functionalization of the PTFE surface with various proteins to test the biological adhesiveness. From the oxidized and crosslinker-coupled film, 12 mm disks were punched and mounted in Minusheet carriers (Minucell) which were placed in 24 microtiter plates. The adhesion protein solution was added and incubated overnight at 4 8C. After being washed with prewarmed PBS, the endothelial cells were left to adhere for 24 h on the PTFE surface under static conditions. For perfusion experiments, six Minusheet carriers were placed in a perfusion container (Minucell) and physiologic shear stress (about 200 s–1) was applied for 72 h, to investigate the stability of the EC lining. The physiological blood viscosity (3.3 mPa s) was mimicked with a DMEM medium (10% FCS, HEPES, PenStrep) containing 6% dextran (70 000 Da).


11.3 Results and Discussion 11.3.1 Wet-Chemical Modification of PTFE Polymer Film

One strategy to reduce occlusion events in small-diameter PTFE grafts can be to line these grafts with endothelial cells. Therefore, a three-step procedure to selectively modify the PTFE polymer surface was developed. Variation in incubation time and temperature changed the gain of reactive hydroxyl groups obtained after the treatment with H2O2/H2SO4. A cross-sectional analysis of the PTFE by SEM showed that the oxidation could be observed to occur in a depth up to 20 nm. Verification of free hydroxyl groups was performed by FTIR. Untreated PTFE showed characteristic peaks between 1000 and 1400 cm–1. After oxidation, new peaks were found at 3300–3500 cm–1 and around 1650 cm–1 (indicative of OH groups). The progression of the oxidation reaction was also followed by measurement of the equilibrium contact angle with water. A significant decrease of the contact angle from 1258 to 88–1158 was noticed, depending on the reaction parameters. No visible changes of the surface morphology were observable by SEM and SFM. While the surface morphology of PTFE remained unchanged after the second reaction step, the reaction could be followed and characterized by FTIR. The spectra showed additional peaks corresponding to C=N bonds at 2200–2400 cm–1 (2450 cm–1). The contact angle with water increased to 1108 after reaction with cyanuric chloride. The stability and elasticity of PTFE remained unaffected. The multiple steps of the modification from oxidation to the attachment of the protein were observed using the contact angle method, but an exact quantification was difficult. The oxidation yielded enough reactive groups such as hydroxyls, ketones, aldehydes, as well as chain breakage. The diameter of the cyanuric chloride molecule is in the region of 1 nm, hydroxyl groups are 0.2 nm and a single protein covers 8–10 nm. For comparison, lipid head groups occupy approximately 0.4 nm2. The hydroxyl group is much smaller and when the protein is covalently bound to the surface, approximately 200 molecules of cyanuric chloride would be covered under the assumption that the surface is homogenous and tightly packed. This indicates that only a low degree of surface activation is required for the intended modification. Changes in surface morphology during modification were followed by SFM. No changes in the overall surface morphology were detectable after oxidation, activation and crosslinking with the protein. This indicates that this method has a good selectivity. Using the high resolution of the AFM technique, it was possible to visualize the attachment of single protein molecules. These appeared as round structures with a diameter of 8–10 nm. This size corresponds well with the molecular weight of Min3 (117 kDa). It was also possible to quantify the degree of modification: There were 408 ± 28 proteins per lm2 on a total analyzed area of 100 lm2 (Fig. 11.4 A). For BSA, collagen IV, and fibronectin a coupling

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Fig. 11.4 AFM images of a Min3-modified PTFE surface (A) before and (B) after the growth of endothelial cells. (A), insert: globular structures with a diameter of 9 nm, comparable in size with the adhesion protein Min3; (B), insert: formation of tight junctions between the individual cells.

density between 95 molecules for collagen and 475 molecules for BSA could be observed. The surface density of biomolecules depends on the size of the molecule, the charge, and the adsorption behavior. The surface modification could be followed further by FTIR spectroscopy. After functionalization, typical protein peaks were detected at 1639 (amide I), 1548, and 1122 cm–1 (amide II), and further peaks at 3415 and 3214 cm–1. 11.3.2 Cell Adhesion Experiments 11.3.2.1 Adhesion and Cultivation in Static Culture The PTFE graft, functionalized with various adhesion molecules, showed an almost confluent lining of the surface with ECs during static culture for all the protein coatings used. Differences could only be observed for the time when maximum confluence was reached (Fig. 11.4 B). Then, only the untreated material and the PTFE coated with BSA (an inactive control) showed an incomplete cell lining. The stability of the EC layers to washing with phosphate buffer increased with rising binding strength of the ligand, in the rank order: BSA < fibronectin < collagen < Min3.

11.3.2.2 Perfusion Experiments Endothelial cells seeded on untreated PTFE or BSA-modified PTFE were totally washed away when shear stress was applied. As shown in Fig. 11.5, only the endothelial cells cultivated on Min3-modified PTFE could withstand the shear


Fig. 11.5 Cell adhesion experiment. Confocal laser scanning microscopy of endothelial cells growing on (A) unmodified and (B) Min3-modified PTFE.

stress, and remained on the surface. Nearly 90% of the total implant surface was still covered with cells after 72 h of perfusion. The ECs were found to be aligned in the flow direction. 11.3.3 Cell Adhesion Force Measurements

In this study, the surface morphology of the implant material and the coupled adhesion molecules could be visualized at molecular resolution in the native wet state by SFM. By scanning the PTFE surface with the cell-modified cantilever, the adhesion properties of the ECs were assessed (Fig. 11.6). When an unmodified or BSA-coated surface was used, a clear image of the probe was achieved. The lateral resolution of the SFM was significantly decreased, because of the decrease in sharpness of the tip on the cantilever (i.e., the ECs were the sensor). When ECM adhesion protein molecules were attached to the surface, strong adhesion of the cantilever/cell nanosensor could be observed, which was evident by the horizontal streaks on the lateral force image (Fig. 11.7 A). This effect could be shown for all adhesion molecules (i.e., Min3, collagen and fibronectin). When the interaction was blocked by injection of an excess of soluble adhesion molecule, a clear image of the surface was regained (Fig. 11.7 B). In force scans (Fig. 11.8), the deflection of the cantilever was recorded on its approach to the PTFE surface and on retraction from the surface, and was directly converted into adhesion force. The calculated results are summarized in Fig. 11.9. Only low adhesion forces were observed by scanning the unmodified or BSA-coated surface (between 0 and 50 pN). When the surface was modified with Min3, fibronectin, or collagen, a drastic increase in the adhesion forces was observed (Fig. 11.9, upper graph). This increase was significantly higher when using Min3, than for fibronectin and collagen. The applied adhesion force

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Fig. 11.6 Experimental set-up for the force measurements. The cantilever with the attached endothelial cells scans the protein-modified PTFE surface and endothelial cells grown on a silicon cantilever (EM image). Insert: 5 lm ´ 7 lm.

Fig. 11.7 Scanning lateral force microscopy. (A) Min3-modified surface; the streaks indicate a strong adhesion during the visualization of the surface; (B) the same surface area was scanned, but the interaction was blocked by soluble Min3.

11.4 Conclusion Fig. 11.8 Individual force vs. distance curves determined by measuring the adhesion of cell-modified cantilevers on modified implant materials. Adhesion time: 2 s.

for all ECM proteins was found to be in the range of some nano-newtons. The total adhesion forces were dependent on the real surface area of the sensor-attached cell, which was in contact with the modified implant material. For each cantilever, varying forces could be observed. The results shown were calculated of the measured forces of at least six individual cantilevers. It could be shown that the adhesion force also depended on the contact time between the cantilever and the implant surface. Longer adhesion times led to higher adhesion forces (Fig. 11.9, lower graph). This fact can be discussed as a time limitation on the bond formation between the individual ligand–receptor pairs and is in agreement with the published literature [38].

11.4 Conclusion

Cell-adsorptive coatings of PTFE with laminin, collagen IV, and fibronectin have been described in the literature to improve the growth of endothelial cells on vascular implants, but these methods did not show satisfactory results when applied to vessels with a small diameter. Here, we report on the covalent attachment of a similar adhesion protein, Min3. The improved performance of this modification can be attributed both to its higher receptor affinity compared with ECM proteins and to the covalent fixation to the PTFE. The latter especially might be of importance, to resist shear stresses as encountered in the systemic circulation. In addition, in this study the force interaction between a human EC and a number of adhesion proteins was assessed for the first time. This was made possible by growing ECs on the tip of an SFM cantilever and scanning across a protein-coated PTFE graft surface. The strongest interaction was found when the bacterial invasin Min3 was used, followed by fibronectin and collagen.

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11 Adhesion Molecule-Modified Cardiovascular Prostheses Fig. 11.9 Upper graph: Calculation of adhesion forces between the protein-modified PTFE surface and the cantilever-attached EC. 1: unmodified surface; 2: BSA; 3: Min3; 4: fibronectin; 5: collagen IV; adhesion time: 2 s. Lower graph: Dependence of calculated adhesion forces between the Min3-modified PTFE surface and the cantileverattached EC on the contact time (between 0.1 and 60 s). The adhesion force increased from 240 pN (0.1 s contact) to 1112 pN (60 s contact).

Acknowledgments

This work was supported financially by the interdisciplinary research program “Biologically Composed Materials and Systems” funded by the Departments of Education, Culture, and Science of the Saarland and Aesculap AG (Tutlingen). We also thank the Fonds der Chemischen Industrie for financial support. U.B. is grateful to the Stiftung Deutscher Naturforscher Leopoldina (BMBF/LPD9901/8-6) for financial support. Special thanks go to Prof. Dr. K. J. Neis and his team from the Frauenklinik of the Caritas Klinik St. Theresia, Saarbrücken, for the supply of umbilical cords.

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tory disease. N. Engl. J. Med., 1999, 340(2): 115–126. Eagle, K. A., Guyton, R. A., Davidoff, R., Ewy, G. A., Fonger, J., Gardner, T. J., Gott, J. P., Herrmann, H. C., Marlow, R. A., Nugent, W. C., O’Connor, G. T., Orszulak, T. A., Rieselbach, R. E., Winters, W. L., Yusuf, S., Gibbons, R. J., Alpert, J. S., Eagle, K. A., Garson, A. Jr., Gregoratos, G., Russell, R. O., Smith, S. C. Jr., ACC/AHA Guidelines for Coronary Artery Bypass Graft Surgery: A Report of the American College of Cardiology/American Heart Association Task Force on Practice Guidelines (Committee to Revise the 1991 Guidelines for Coronary Artery Bypass Graft Surgery). American College of Cardiology/American Heart Association. J. Am. Coll. Cardiol., 1999, 34(4): 1262–1347. Okies, J. E., Page, U. S., Bigelow, J. C., Krause, A. H., Salomon, N. W., The left internal mammary artery: the graft of choice. Circulation, 1984, 70(3 Pt 2): I213–I221. Catinella, F. P., Cunningham, J. N. Jr., Srungaram, R. K., Baumann, F. G., Nathan, I. M., Glassman, E. A., Knopp, E. A., Spencer, F. C., The factors influencing early patency of coronary artery bypass vein grafts: correlation of angiographic and ultrastructural findings. J. Thorac. Cardiovasc. Surg., 1982, 83(5): 686–700. Lytle, B. W., Loop, F. D., Taylor, P. C., Simpfendorfer, C., Kramer, J. R., Ratliff, N. B., Goormastic, M., Cosgrove, D. M., Vein graft disease: the clinical impact of stenoses in saphenous vein bypass grafts to coronary arteries. J. Thorac. Cardiovasc. Surg., 1992, 103(5): 831–840. Hunter, G. C., Carson, S. N., Wong, H. N., French, S., Experimental small-diameter graft patency: effect of compliance, porosity and graft healing potential. Curr. Surg., 1980, 37(6): 439–441. Herzberg, M. C., Coagulation and thrombosis in cardiovascular disease: plausible contributions of infectious agents. Ann. Periodontol., 2001, 6(1): 16–19.

8 Smith, E. B., Haemostatic factors and

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atherogenesis. Atherosclerosis, 1996, 124(2): 137–143. Smith, E. B., Fibrinogen, fibrin and fibrin degradation products in relation to atherosclerosis. Clin. Haematol., 1986, 15(2): 355–370. Balkuv-Ulutin, S., Fibrinolytic system in atherosclerosis. Semin. Thromb. Hemost., 1986, 12(2): 91–101. Thomson, G. J., Vohra, R. K., Carr, M. H., Walker, M. G., Adult human endothelial cell seeding using expanded polytetrafluoroethylene vascular grafts: a comparison of four substrates. Surgery, 1991, 109(1): 20– 27. Deutsch, M., Eberl, T., Fischlein, T., Meinhart, J., Minar, E., Puschmann, R., Schmid, P., Zilla, P., In vitro endothelialization of ePTFE vascular prostheses in clinical use: preliminary results. Vasa Suppl., 1990, 30: 219–220. Deutsch, M., Meinhart, J., Vesely, M., Fischlein, T., Groscurth, P., von Oppell, U., Zilla, P., In vitro endothelialization of expanded polytetrafluoroethylene grafts: a clinical case report after 41 months of implantation. J. Vasc. Surg., 1997, 25(4): 757–763. Kadletz, M., Magometschnigg, H., Minar, E., Konig, G., Grabenwoger, M., Grimm, M., Wolner, E., Implantation of in vitro endothelialized polytetrafluoroethylene grafts in human beings. A preliminary report. J. Thorac. Cardiovasc. Surg., 1992, 104(3): 736–742. Meinhart, J., M. Deutsch, P. Zilla, Eight years of clinical endothelial cell transplantation. Closing the gap between prosthetic grafts and vein grafts. ASAIO J., 1997, 43(5): M515–M521. Zilla, P., Deutsch, M., Meinhart, J., Puschmann, R., Eberl, T., Minar, E., Dudczak, R., Lugmaier, H., Schmidt, P., Noszian, I., Clinical in vitro endothelialization of femoropopliteal bypass grafts: an actuarial follow-up over three years. J. Vasc. Surg., 1994, 19(3): 540–548. Pasquinelli, G., Freyrie, A., Preda, P., Curti, T., D’Addato, M., Laschi, R., Heal-

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ing of prosthetic arterial grafts. Scanning Microsc., 1990, 4(2): 351–362. Marois, Y., M. F. Sigot Luizard, R. Guidoin, Endothelial cell behavior on vascular prosthetic grafts: effect of polymer chemistry, surface structure, and surface treatment. ASAIO J., 1999, 45(4): 272–280. Zilla, P., Preiss, P., Groscurth, P., Rosemeier, F., Deutsch, M., Odell, J., Heidinger, C., Fasol, R., von Oppell, U., In vitro-lined endothelium: initial integrity and ultrastructural events. Surgery, 1994, 116(3): 524–534. Sipehia, R., G. Martucci, J. Lipscombe, Transplantation of human endothelial cell monolayer on artificial vascular prosthesis: the effect of growth-support surface chemistry, cell seeding density, ECM protein coating, and growth factors. Artif. Cells Blood Substit. Immobil. Biotechnol., 1996, 24(1): 51–63. Conte, M. S., Choudhury, R. P., Shirakowa, M., Fallon, J. T., Birinyi, L. K., Choudhry, R. P., Endothelial cell seeding fails to attenuate intimal thickening in balloon-injured rabbit arteries. J. Vasc. Surg., 1995, 21(3): 413–421. van Wachem, P. B., Vreriks, C. M., Beugeling, T., Feijen, J., Bantjes, A., Detmers, J. P., van Aken, W. G., The influence of protein adsorption on interactions of cultured human endothelial cells with polymers. J. Biomed. Mater. Res., 1987, 21(6): 701–718. Radomski, J. S., Jarrell, B. E., Williams, S. K., Koolpe, E. A., Greener, D. A., Carabasi, R. A., Initial adherence of human capillary endothelial cells to Dacron. J. Surg. Res., 1987, 42(2): 133–140. Schmidt, S. P., Hunter, T. J., Sharp, W. V., Malindzak, G. S., Evancho, M. M., Endothelial cell-seeded four-millimeter Dacron vascular grafts: effects of blood flow manipulation through the grafts. J. Vasc. Surg., 1984, 1(3): 434–441. Ramalanjaona, G., Kempczinski, R. F., Rosenman, J. E., Douville, E. C., Silberstein, E. B., The effect of fibronectin coating on endothelial cell kinetics in polytetrafluoroethylene grafts. J. Vasc. Surg., 1986. 3(2): 264–272. Kadletz, M., Moser, R., Preiss, P., Deutsch, M., Zilla, P., Fasol, R., In vitro

27

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lining of fibronectin coated PTFE grafts with cryopreserved saphenous vein endothelial cells. Thorac. Cardiovasc. Surg., 1987, 35 (Spec. No. 2): 143–147. Walluscheck, K. P., Steinhoff, G., Kelm, S., Haverich, A., Improved endothelial cell attachment on ePTFE vascular grafts pretreated with synthetic RGD-containing peptides. Eur. J. Vasc. Endovasc. Surg., 1996, 12(3): 321–330. Chan, C., T. Ko, H. Hiraoka, Polymer surface modification by plasmas and photons. Surf. Sci. Rep., 1996, 24: 1. Niino, H., A. Yabe, Surface modification and metallization of fluorocarbon polymers by excimer laser processing. Appl. Phys. Lett., 1993, 63: 3527. Niino, H., A. Yabe, Chemical surface modification of fluorocarbon polymers by excimer laser processing. Appl. Surf. Sci., 1996, 96–98: 550. Beyer, D., Knoll, W., Ringsdorf, H., Wang, J. H., Timmons, R. B., Sluka, P., Reduced protein adsorption on plastics via direct plasma deposition of triethylene glycol monoallyl ether. J. Biomed. Mater. Res., 1997, 36(2): 181–189. Binnig, G., C. F. Quate, C. Gerber, Phys. Rev. Lett., 1986, 56: 930–933. Hansma, H. G., K. J. Kim, D. E. Laney, R. A. Garcia, M. Argaman, M. J. Allen, S. M. Parsons, Properties of biomolecules measured from atomic force microscope images: a review. J. Struct. Biol., 1997, 119, 99–108. Radmacher, M., J. P. Cleveland, M. Fritz, H. G. Hansma, P. K. Hansma, Mapping interaction forces with the atomic force microscope. Biophys. J., 1994, 66(6): 2159– 2165. Weisenhorn, A. L., M. Egger, F. Ohnesorge, S. A. C. Gould, S.-P. Heyn, H. G. Hansma, R. L. Sinsheimer, H. E. Gaub, P. K. Hansma, Molecular-resolution images of Langmuir-Blodgett films and DNA by atomic force microscopy. Langmuir, 1991, 7: 8–12. Löhbach, C., U. Bakowsky, C. Kneuer, D. Jahn, T. Graeter, H.-J. Schäfers, C.-M. Lehr, Wet chemical modification of PTFE implant surfaces with a specific cell adhesion molecule. Chem. Commun., 2002, 21: 2568–2569.

References 37 Florin, E. L., Moy V. T., Gaub, H. E.,

Science 1994, 264, 415–417. 38 Henderson, R., H. Oberleithner, Pushing, pulling, dragging, and vibrating renal

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12 Surface Engineering by Coating of Hydrophilic Layers: Bioadhesion and Biocontamination G. Legeay and F. Poncin-Epaillard

Abstract

Plasma techniques – plasma modification, plasma deposition and plasma and grafting reactions – are illustrated for the elaboration of tools, medical devices and biomaterials. These promising techniques are developed either for the sterilization, antifouling surfaces or for enhancing the biofilms formation.

12.1 Introduction

Biomaterial corresponds to any type of material (e.g., natural, synthetic, or part of a natural structure) that increases or replaces a natural function. Many types of materials are used, including polymers, metals, ceramics, glasses, composites, and alloys, in devices substituting for heart valves, joints, tubing, intraocular lenses, etc. The requirements are compatibility with living tissues (minimal property alterations, minimal local deleterious effects, absence of harmful systemic effects), appropriate mechanical properties (elasticity, yield stress, ductility, toughness, deformation, fatigue, hardness, wear resistance), and suitable manufacturing methods, quality, sterilization techniques, cost and so on). Their surface chemistry and their structural properties largely control the biological response. 12.1.1 The Need for Bioadhesion of Biomaterials

Bioadhesion, i.e. biofilm formation resulting in a fouling surface, is required for biomaterial to be considered as a part of the body (e.g., orthopedic prosthesis, hard tissue) to enhance its incorporation and its biomechanical response. Examples are in the rebuilding of bones, recolonization, and hybrid implants composed of two parts, a synthetic one (with polymers as the mechanical subAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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12 Surface Engineering by Coating of Hydrophilic Layers: Bioadhesion and Biocontamination

strate), and an animal or human one (such as epithelial cells). Biofilm formation that induces a fouling surface is not necessary with temporary materials like tubes, drains, patches, pumps or dental prostheses; with biomaterials bearing specific characteristics (optical or permeation properties); with materials implanted in an area with high mechanical strength (friction); or with medico-surgical tools avoiding any bacterial transmission (nosocomial diseases, human immunodeficiency virus (HIV) titration, prion proteins). The nature of the interactions between the biological surroundings and the biomaterial at the interface, and the following biological response, depend on the material surface properties such as surface energy, chemical composition, cleanliness, texture, corrosion resistance, and alteration of neighboring proteins on the micrometer and nanometer scales. To avoid biomaterial synthesis by a trial-or-error method, control of (non)bioadhesion related to easier control of surface properties must be achieved. Therefore, the surface energy and the work of adhesion of the biosurface with a model protein, bovine serum albumin (BSA), could be interpreted as pertinent parameters [1]. The interactions between the surface and the protein, and the resulting protein adhesion, are the highest when the material surface energy is around 50 mJ m–2, corresponding to a hydrophilic/hydrophobic equilibrium balance, and the lowest for either polar or apolar surfaces. Although this value helps to characterize the possibility of an available biosurface with high or low energy, it could not be the absolute criterion since it has been developed for only a single protein. 12.1.2 Mechanism of Bioadhesion

The bioadhesion mechanism is not well defined. The literature indicates that bioadhesion takes place in three steps: the deposition and attachment of a protein layer, followed by cell deposition. The important parameter corresponds to the formation of a water monolayer on the substrate (Fig. 12.1). Because of its low wettability (and high water contact angle), a substrate with a low surface energy could interact neither with biomolecules nor with water molecules. There is no affinity, only electrostatic repulsion of the proteins. Hence, the organization of the biomolecules at the interface is not uniform, and cell aggregates are formed. Conversely, with high surface energy materials, a low-cohesion (no Young’s modulus) monomolecular layer of water is formed spontaneously and characterized by several features: the water tension is higher than that of a protein, and the hydrophilic surface affinity with the small water molecules is stronger than with more complex molecules. Therefore, a front of small and mobile water molecules is created during wetting and leads to water spreading on the surface, and monolayer formation. Sometimes, interface swelling is observed within the first few seconds. Hydrogen bonds are established immediately between the polar groups of the polymer and the water molecules. Salt ions (Na+, Cl–) enhance the phenomenon. Then, the anchoring of hydrophilic biomolecules

12.2 Surface Engineering

Fig. 12.1 Adhesion behavior of hydrophobic and hydrophilic materials.

(protein derivatives, i.e., BSA) by hydrogen bonds takes place, but without any adhesion to the substrate. With such materials, the interface is defined by the aqueous top layer of substrate and by the surface of the cells or bacteria, with a protein layer in between. In conclusion, the deposition of a water monolayer on a hydrophilic substrate prohibits the formation of a stable biomolecular layer and cells cannot be deposited.

12.2 Surface Engineering

Surface engineering means both surface preparation and sterilization of the modified surface. A recent review of surface engineering [2] indicates that plasma surface preparation is well adapted to the preparation of biomaterials because it is easy to reproduce and clean the treated sample, and good for the environment (no pollution) as well. This type of treatment can also be applied in a clean room. Several plasma processes are possible for surface engineering. 12.2.1 Surface Preparation

The plasma technique offers the possibility of changing the chemical composition and properties of a surface continuously, and thus the biocompatibility of the material. The plasma is mostly created by an electrical glow discharge in vacuum. So the reactive species (ions and radicals) can interact with the whole surface. Nowadays it appears as a versatile process in biomaterials engineering, because it has several advantages. It can be applied to any type of sample geometry. Hydrophilic or hydrophobic surfaces can be obtained. The thickness of the modified layer depends on the type of process (simple surface functionalization, deposition, or postgrafting). Surface functionalization modifies the material to a depth of about 10 nm, without any alterations of mechanical or optical properties. Plasmas like oxygen or carbon dioxide induce the attachment of polar groups (–C–OH, –C=O, –COOH, etc.).

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12.2.2 Surface Sterilization

Before their use or any biological testing (in vitro and in vivo), sterilization of biomaterials and medical devices is necessary. The sterilization process must not cause damage or chemical alterations to the materials. Physical or chemical techniques have been developed, some of which are illustrated here. Water steam sterilization at 121 8C for 20 min is a popular technique but polymers are always degraded and hydrophilic surfaces can swell. Irradiation of polymers with a c beam leads to C–C and C–O bond scissions and the formation of unstable radicals. After the recombination of these highly reactive species, oxidation and crosslinking are observed. Exposure of materials to ethylene oxide at 40–45 8C for about an hour leads to gas desorption lasting from several hours to several days. However, no chemical damage is observed. This technique is preferable for hydrophilic materials. Whatever the sterilization process is, it must be integrated into any concept for new materials and surface engineering.

12.3 Results and Discussion 12.3.1 Hydrophobic Cold Plasma Treated Surfaces in Ophthalmology

Most plasma-treated hydrophobic surfaces of biomaterials are formed with tetrafluoromethane (CF4) plasma interactions [3, 4]. The modified surface represents a nonadherent polytetrafluoroethylene (PTFE)-like structure (–(CF2)n–) with low surface energy that could vary from 20 mJ m–2 down to only a few mJ m–2 when the super-hydrophobic character is pronounced. The chosen operating conditions lead to a low fluorine atom density in the plasma, thus avoiding surface degradation. Such surfaces are applied in order to prevent the formation of the biofilm. Poly(methyl methacrylate) (PMMA) intra-ocular lenses provide an example. The adhesion of proteins, the resulting development of inflammatory cells, and the formation of cellular debris are mostly avoided when the substrate is treated in CF4 plasma. The modified surface corresponds to a smooth PTFE-like structure and the initial optical properties are not altered. As a result, cell activation and granulocyte adhesion are reduced.


12.3.2 Hydrophilic Cold Plasma Treated Surfaces Based on Polyvinylpyrrolidone (PVP) or Natural Derivative Coatings

As discussed above, the synthesis of either a hydrophobic or a hydrophilic surface avoids biofilm formation. In case of a hydrophilic surface, the different techniques – modification, deposition, and post-grafting – can be explored, but here only post-grafting will be described because this type of modification results in an improved control of aging. Hydrophilic monomers or polymers were grafted onto a plasma-pretreated surface. The hydrophilicity of several hydrophilic polymers was measured. The corresponding contact angles and the calculated surface energy values are given in Table 12.1. PVP possessed the highest surface energy and also the most polar component. Its grafting onto different substrates led to a surface with a surface energy (cs) of 63.4 mJ m–2 and a polar component (cp) of 42.0 mJ m–2. The acid–base characteristics of other polymers, such as the cellulose derivatives hydroxyethylcellulose (HEC) and carboxymethylcellulose (CMC), should be interesting. These hydrophilic polymers bear carboxylic groups that are able to attract the amino groups of proteins and the surface proteins of bacteria and cells.

Table 12.1 Contact angle (sessile drop) and surface energy of several hydrophilic polymers (calculated by the Owens Wendt method).

Polymer a

HPMC (E4M) CMC HEC HPC EC PVP PVA (75% OH) PVA (88% OH) PVA (95% OH) Poly(HEMA) a)

Contact angle [8] water

formamide

CH2I2

60.4 44.0 28.4 61.6 73.7 26.0 49.0 52.9 76.4 42.3

55.3 19.3 34.1 49.7 67.6 34.9 34.2 40.0 40.2 16.3

49.7 48.2 34.1 43.1 50.6 45.2 36.2 35.2 46.4 35.2

HPMC: hydroxypropylmethylcellulose; CMC: carboxymethylcellulose; HEC: hydroxyethylcellulose; HPC: hydroxypropylcellulose; EC: ethylcellulose; PVP: polyvinylpyrrolidone; PVA: poly(vinyl alcohol); poly(HEMA): poly(hydroxyethyl methacrylate).

cs [mJ m–2]

cp [mJ m–2]

42.1 55.9 62.5 43.7 34.0 63.4 52.9 50.3 41.6 58.2

18.5 29.5 36.9 15.7 9.9 42.0 22.8 19.9 6.5 27.0

179

180


12.3.2.1 Grafting of Monomer onto Plasma-Pretreated Surfaces The post-grafting reaction provides a strong and irreversible adhesion between the grafted layer and the substrate, since a covalent bond is established. We grafted N-vinylpyrrolidone (NVP) monomer onto a polypropylene (PP) surface pretreated with cold N2 plasma (Scheme 12.1). After the plasma treatment, the modified surface was dipped in aqueous NVP solution for several periods of time (1–50 h) and at various temperatures (60–80 8C). Plasma treatment of PP induces the formation of either functional groups (mostly amino groups) or radicals whose concentration depends on plasma parameters. The synthesis route is divided in two parts: 1. Activation of surfaces to create radicals: under cold plasma conditions, C–C, C–O, and C–H bonds break and the resulting free radicals appear in a surface layer about 10 nm thick. These unstable radicals may react through a radical mechanism with a monomer such as NVP unless, after the plasma treatment, the substrate is immediately dipped into the monomer solution under inert atmosphere. 2. Polymerization: heating of this solution with the plasma-treated substrate causes new macromolecular chains to grow onto the surface. The alkyl or peroxide radicals, whose chemical nature depends on the atmosphere after the plasma treatment, initiate the grafting reaction of NVP monomer onto the polymer surface. The grafting rate of the polymer chains increases with the surface radical density, which depends on the plasma parameters. The grafted NVP chain density is obtained from a conductimetric back-titration of carboxylic groups derived from NVP ring opening. The grafting yield also depends on the grafting time (Fig. 12.2 a). When initiated with alkyl radicals (PP is kept under a N2 atmosphere after plasma activation), the grafting is rapid during the first 10 min. Later the grafted polymer increases more slowly as the reaction pro-

Scheme 12.1 Post-grafting of NVP onto plasma-treated PP.


Fig. 12.2 Variation with grafting time and the initiator type of (a) grafting yield; (b) length of the grafted PVP chains.

gresses further. At the beginning the reaction is initiated by more reactive radicals at the upper layer. When grafting is initiated by hydroperoxides, the kinetics is slower, mainly because the initiation step induces hydroperoxide decomposition and then the reaction between an alkoxy radical and a monomer molecule. The grafted chain length corresponds to almost 20 monomer units for alkyl radicals and 12 for alkoxy radicals. The alkyl radical produces longer chains (Fig. 12.2 b), calculated by comparison of the density of titrated radicals with the amount of grafted NVP (molecular weight of grafted chain: about 2000 Da). The chemically grafted PVP chains are shorter than those only deposited on the substrate, as described in Section 12.3.2.2, but long enough for biological recognition that takes place across a distance of 1–2 nm. The comparison of the O/C and N/C element ratios determined by X-ray photoelectron spectroscopy (XPS) analysis (Table 12.2) and the theoretical values suggests that the grafting of NVP covers the PP surface only partially since the N : C and O : C ratios are slightly lower than for PVP. In addition, the C–C/C–H peak is higher than for PVP, indicating that the average thickness of the grafted PVP is less than the XPS information depth (£ 10 nm).

12.3.2.2 Coating with Commercial Native or Synthetic Polymers The coating consists of polyvinylpyrrolidone or cellulosic derivatives deposited by dipping the polymeric substrate into the diluted polymer solution (Scheme 12.2). Since most of the usual polymeric substrates (acrylic resin, polyethylene, etc.) are hydrophobic, they must be pretreated before the coating process. The polymeric substrate surface have to be more polar than the aqueous solutions. Increasing their surface energy implies the formation of a uniform coating. As

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12 Surface Engineering by Coating of Hydrophilic Layers: Bioadhesion and Biocontamination Table 12.2 XPS data of PP-grafted NVP.

Sample

O/C (´ 100)

N/C (´ 100)

Virgin PP N2 plasma-treated PP Plasma-treated PP grafted with NVP PVP (theoretical values)

5.3 14.3 10.9

0 28.6 10.3

16.5

16.5

Assignment

Relative intensity [%]

Theoretical intensity [%]

C–C, C–H C–N N–C=O

61 26 13

51 33 16

Chemical structure of PVP-PP

the substrate has been chosen for its bulk properties (mechanical, etc.) the pretreatment must not alter these properties. Therefore, cold Ar or O2 plasma is applied. It is produced by RF (13.56 MHz) or microwave (254 MHz) discharges. As before, the synthesis is also divided in two parts: 1. Surface activation to create radicals by the same scheme as described above: After the plasma treatment, the substrate is kept in an air atmosphere and most of the radicals formed are oxidized to give oxygen functionalities (–C– OH, –C=O, etc.). As a result of the activation, the surface energy of the substrate increases and allows the uniform deposition of the coating. 2. Coating with native and synthetic polymers: This simply involves dipping the substrate into the aqueous solution of the native or synthetic polymer and subsequent drying. Finally, the polymer film coats all the substrate to a thickness that depends on the dilution. It appears that the coating resists any removal operation such as Soxhlet extraction more strongly than a simple deposition without any pre-activation. Some radicals formed during the plasma treatment would probably react with the polymer film. Surface examination by XPS analysis [5] before and after washing with water (Table 12.3) and even with phosphate detergent (e.g., with one of the solutions usual in hospitals, such as RBS® detergent, consisting of a mixture of

Scheme 12.2 Plasma activation of the substrate, and polymer deposition.


183

Table 12.3 Elemental composition (XPS analysis) of different membranes.

C1s [%]

O1s [%]

Binding energy [eV]

284.70

Assignment

C–C, C–H C–N

286.2

PC reference

60.84

12.75

Membrane as received

54.30

H2O/EtOH-washed

60.60

Ar plasma + PVP (1%) Ar plasma + PVP (1%) + H2O rinse PVP reference a)

287.70

290.44

291.9

534

532.3

O=C–N

O–COO–

a)

C–O–C

C=O

531.3 N–C=O

–

6.69

3.61

10.77

5.35

–

18.09

3.94

3.72

0.87

7.65

5.68

2.74

15.15

–

4.27

1.87

9.13

5.99

1.60

46.39

21.14

7.52

2.53

0.56

4.67

6.42

5.38

48.75

21.58

5.66

1.88

0.14

5.54

7.10

4.78

33.49

33.57

9.44

0.00

0.00

0.00

1.40

11.01

p–p* transition peak.

anionic and nonionic surfactants) indicates the presence of PVP coating on the substrate after the washing. These results are also found by ToF-SIMS spectroscopy [6], and lead to the idea of the formation of a thin polymer film covering the surface. The adhesion of the thin layer is obtained by hydrogen bonding which resists the washing. Rinsing of biomaterials is also one of the possible ways to eliminate the biofilm, but after its deposition, not during its formation, except in dynamic systems. For several years, native biosurfactants have been being proposed and used as washing agents. For example, nonionic lipopolysaccharide biosurfactants reduce the bioadhesion of bacteria [7–9]. They are obtained by biotechnology, specifically as secretions from specific cells. These lipopolysaccharides (C12E5, C12E7, C12E9, C12E10) possess a surface energy (c) of about 36 mJ m–2, which is comparable with the value for Escherichia coli (about 38.4 mJ m–2). The biosurfactants seem to be more efficient than the synthetic surfactants (pentaethylene glycol monododecyl ether) in reducing the bioadhesion of E. coli K12, i.e., fewer bacteria adhere to the substrate cleaned with the biosurfactant. These remaining bacteria are counted by epifluorescence microscopy. To avoid the supplementary step of washing, hydrophilic coatings bearing drugs, antibiotics or antistenosis agents have been synthesized [10]. The destruction of the biofilm was shown to depend on the kinetics of release of the drug into the biological environment. 12.3.3 Examples

When biological tests are applied, all the samples are prepared in a clean room (class 100).

184


12.3.3.1 With Different Biomolecules, i.e., Proteins On a native (hydrophobic) polycarbonate membrane, the observed albumin adsorption is dominated by electrostatic interactions. No bioadhesion of bacteria is required in order to protect the permeation properties of the membrane. The evaluation of a polycarbonate membrane coated with polyvinylpyrrolidone indicates that it has a lower affinity for immunoglobulins (IgG) [5] and for albumin [6] than the native one. Hence, the bioadhesion is lower. The protein conformation can be modified by a surface effect related to the hydrophilicity/hydrophobicity and also to the electronegative/electropositive character. The hydrophilic membrane (polycarbonate/PVP) surface modifies the molecular conformation of albumin, increases its water solubility, and as a result decreases its adsorption onto the membrane.

12.3.3.2 Implantation (ex in vivo) The plasma activation also improves the adhesion of polymer films applicable in odontology for dental prostheses (ex in vivo implantation) [11–15]. Completely immovable dental prostheses composed of steel or acrylic polymers exhibit bad contact with the palate, proliferation of cells and formation of a biofilm, friction, and irritation of tissue. Adhesion of a dental prosthesis is enhanced by the deposition of a specific glue, polyvinylpyrrolidone, or cellulose derivatives on the plasma-treated prosthesis [16–19]. The prosthesis is treated in Ar plasma to increase its wettability and surface energy before it is coated with a 10 lm thick hydrophilic polymer (either PVP or cellulose derivatives). Such a biomaterial provides a good cohesion, a low cellular proliferation as shown with in-vitro tests in synthetic saliva medium, good mechanical fixation to the palate without dental glue, and a good in-vivo stability [11–13]. Patients also have a subjectively good appreciation associated with correct contact with the palate, better elocation, and a better feel. Such processes are also applied to plastic reconstruction of mammalian prostheses [20] and tubing [21]. A polyvinylpyrrolidone coating was also applied to polyethylene pretreated in a cold O2 plasma (RF, 20 L, 50 W, 10 min). The plasma treatment induces the attachment of alcohol and ketone groups, and therefore a hydrophilic character develops. Then, the plasma-modified polyethylene substrate is dipped into an aqueous solution of PVP (Kollidon 90 pharmaceutical grade, molecular weight: 90 000 Da). A new, stable, hydrophilic substrate is obtained. Polystyrene treated with Ar plasma (with Petri dishes as substrates) was also grafted with cellulose derivative polymers [22], then mouse melanoma cells (B16, C3, F10) were spread onto the treated surfaces. All three cell lines proliferated significantly more slowly after 48 h [23]. Further, we showed that cells aggregated on this modified substrate could rapidly establish gap junctions for communication between the fibronectin of cell membranes and exogenic surfaces [24]. Cell differentiation occurs and can be used in the biological analysis. For example (Fig. 12.3), pyramidal aggregates of melanoma cells (Swiss 3T3) appear on polystyrene substrates coated with cellulose derivatives after 24 h. These

12.3 Results and Discussion Fig. 12.3 Pyramidal aggregate of melanoma cells (Swiss 3T3) after 24 h on a coating of cellulose derivatives on a polystyrene substrate. The differences in radius of the concentric regions correspond approximately to the size of a cell.

50 lm

pyramids correspond to the stacking of concentric layers whose thickness is approximately equal to the cell size. After several days apoptosis of the cells appears.

12.3.3.3 In vivo Implantation With native urine as a fluid, E. coli and other bacteria can be deposited and anchored onto the substrate. After their growth, multiplication and dissemination are observed. The results are an inflammatory reaction and, sometimes, the appearance of a new illness of the patient. Evaluation of bioadhesion of, and biocontamination by, E. coli on a urinal catheter coated with PVP indicates a decrease in deposited bacteria as compared with the virgin material (Fig. 12.4). The substrate friction coefficient is also decreased and that will ease the introduction of the prosthesis (plastic reconstitution, tubing, etc.) into the organ. On urological devices, the heparin hydrophilic [25] or Lactobacillus fermentum protein [26] coatings were recently proposed to prevent bacterial adhesion on catheters; illustrations of the proliferation kinetics on tubing are given with two common uropathogenic species (E. coli and Enterococcus faecalis). Inclusion of antibiotics in the hydrophilic coating was also proposed [25]. Upon urine flow, antibiotics become free and attack biofilm formation. But the antibiotics may attack the endothelial cells of the patient and cause health problems. New cardiovascular applications are now appearing to control the surface properties of stents and tubing (catheters for stenosis proliferation) [27, 28]. Stents based on titanium or steel are pretreated by dipping them into acid solution and subsequently coating them with a drug (heparin, paclitaxel, etc.). In a biological environment, the deposited layer swells and the drug is released during a period of several days.

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40 lm __

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Fig. 12.4 Colonization by E. coli K12 (HH97531012) after 24 h of dipping in urine: (a) untreated sample; (b) sample coated with PVP.

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Fig. 12.5 Sample of a polycarbonate film after one month of implantation in the peritoneal cavity of a rat: (a) normal membrane; (b) membrane with a PVP coating.

The hydrophilic coating of polyvinylpyrrolidone decreases the bioadhesion of endothelial cells [5] (Fig. 12.5) during implantation in the peritoneal cavity of the rat. After one month, the coated polycarbonate (PC) surface shows a cellular density of less than 5% of the surface. Without the hydrophilic coating, cells cover about 75% of virgin PC surface. The PC results in a good biocompatibility, no inflammation, and no thrombogenic properties [29].

12.4 Conclusion

Biocompatibility is shown to be correlated to the surface properties and to the possibility of deposition and attachment of a protein layer, i.e., a biofilm. The

References

substrate with low surface energy could interact neither with the biomolecule nor with water molecules. With a material of high surface energy, a water monolayer forms spontaneously. Many materials are too hydrophilic or too hydrophobic to promote the biofilm growth. Therefore, the plasma techniques provide ways of surface engineering a material to provide cell adhesion. The treated surfaces possess an altered wetting behavior and improved cell adhesion. Contrary to classical chemical routes, all these plasma techniques are dry processes and more reliable for biomedical applications. In cases where bioadhesion has to be avoided, such as on medical devices, in hygienic prevention (e.g., nosocomial illness), or in food packaging, the use of substrates with a very hydrophilic character may help to eliminate the proliferation of a biofilm of cells and bacteria. Such a technique is so efficient that antibiotic molecules are not necessary. Therefore, surface engineering is expected to grow rapidly in the field of modifying and adapting materials for specific biological environments. The number of relevant patents and papers is increasing strongly and some of this work has already led to new, commercial, well adapted biomaterials. These applications focus on prostheses, supports for biological assays, and also on all the medical devices used in surgery and in hospital. The same approach could also be adopted for packaging in the food industry.

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M. N. Bellon-Fontaine, M. Hegraud, Appl. Environ. Microbiol., 2002, 68, 728– 737. G. Chang, J. Adhesion Sci. Technol., 2003, 17(16), 2131–2139. C. Wright, G. H. Llanos, R. Rakos, K. King, US Patent 017 69 15, 2003. G. Legeay, F. Poncin-Epaillard, S. Bellessort, Proc. CIP, 2001, 2001, 10–11. S. Suzer, N. Ozden, F. Akaltan, G. Akovali, Appl. Spectrosc., 1997, 51, 1741– 1749. G. Legeay, S. Bellessort, European Patent 1 133 972, 2001. F. Hostettler, M. N. Helmus, Ni Ding, US Patent 5 849 368, 1997. S. Masakazu, European Patent 59 164 709, 1984. J. E. Grasso, Letters to the editor – denture adhesives. JADA, 1996, 124–127. S. Bellesort, Stratégie prothétique 2002, 2(2), 113–122.

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Williams, D. Stickler, E. Nagy, Int. J. Antimicrobial Agents, 2004, 23, S67–S74. P. Cadieux, J. D. Watterson, J. Denstedt, R. R. Harbottle, J. Puskas, J. Howard, B. S. Gan, G. Reid, Coll. Surf. B: Biointerfaces, 2003, 28, 95–105. G. Kopia, G. Llanos, R. Falotico, WO 01/ 87372/A1, 2001. D. Fischell, J. Spaltro, WO 2004/ 002547 A2, 2004. G. Legeay, P. Bertrand, A. Belcourt, L. Kessler, PCT/FR02/00347, 2002.

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13 New Resins and Nanosystems for High-Performance Adhesives R. Mülhaupt

Abstract

Recent advances in polymer sciences and nanotechnology are stimulating the development of adhesive with unusual property profiles. New catalyst systems and controlled polymerization processes afford unprecedented control of molecular and supermolecular architectures. Single site catalyst technology is the key to tailor-made polyolefins ranging from functional waxes and tackifiers to hot melt adhesives. Biocatalysis is claiming an increasing role in production of functional polymers. Advanced reactive extrusion processes exploit novel chain extender chemistry to produce block and graft copolymers as well as functionalized monomers and particles in solvent-free polymer extrusion processes. Carbonylbiscaprolactam is a new non-toxic additive which converts amines and hydroxyl groups into caprolactam blocked isocyanates without requiring the use of isocyanates or phosgene (“isocyanate-free polyurethane chemistry”). Nanophase separation and nanoparticle assembly improve the performance of flexible and structural adhesives. Epoxy adhesives based upon nanophase separated liquid rubber blends afford extraordinary crash resistance. The assembly of nanometer-scaled hyperbranched and dendritic polymers in bulk and at surfaces promotes both adhesion and cohesion. New families of nanocomposites based upon the assembly of organophilic nanoparticles such as silica, boehmite and layered silicates exhibit skeleton-like superstructures and afford the attractive combination of matrix reinforcement, high dimensional and environmental stabilities, improved toughness/stiffness balance, barrier resistance, halogen-free fire protection, tear resistance, and surface adhesion. Progress is highlighted by means of selected examples of new polymer intermediates and nanostructure formation.


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13 New Resins and Nanosystems for High-Performance Adhesives

13.1 Introduction

Joining together a great variety of materials, such as metal, ceramics, glass, wood, rubber, paper, fibers, thermoplastics, and thermosets, represents a key challenge in modern technologies, from automotive and aerospace applications to the assembly of energy conversion systems [1]. The uniform stress distribution in bond lines and the use of advanced formulated polymer systems offers very attractive benefits relating to cost and weight saving combined with durability and corrosion resistance. According to the late Ralph Drake, the development of structural adhesives has seen two decades of enduring progress [2]. Adhesives development requires intensive collaboration in multidisciplinary fields including polymer sciences, polymer processing, engineering, materials, nano, and surface sciences. As illustrated in Fig. 13.1, it is important to balance the requirements for good cohesive and adhesive strengths without sacrificing easy processing and durability, in order to match the demands of market needs and those of sustainable development. This overview highlights recent advances in the development of new polymeric materials and nanosystems for adhesive applications.

13.2 Tailor-Made Polymers and Properties on Demand

Since the mid-1980s significant progress has been made with respect to the precision of polymer synthesis ranging from free-radical polymerization to advanced olefin polymerization and biotechnology. Novel structural and functional polymers are now at hand and offer new opportunities for improving adhesive formulations. Modern polymers can be tailored to offer problem solutions and properties on demand. This includes molecular design in conjunction with con-

Fig. 13.1 Adhesive development.


trolled superstructure formation and formulation of multicomponent multiphase systems. 13.2.1 Controlled Polymerization and Catalysis

Precise control of molecular architectures has been achieved by means of highly stereoselective polymerization processes, excellent control of comonomer incorporation, and end-group variation. During the 1950s living polymerization and the related block copolymer syntheses were restricted to anionic styrene homoand copolymerization. Today, controlled polymerization processes are available to produce a very wide range of block copolymers by means of free-radical, cationic, anionic, and insertion-type polymerization involving most of the cheap monomer feedstocks from petrochemistry. A characteristic feature of controlled radical polymerization, also known as atom transfer polymerization, inifer technology, tetramethyl-N-oxypiperdine (TEMPO)-mediated polymerization, or radical addition/fragmentation transfer (RAFT) polymerization, is the presence of a very rapid equilibrium between active and temporarily inactive (“dormant”) chain ends. This reversible deactivation prevents irreversible chain termination. The remarkable progress of controlled polymerization has expanded the scope of free-radical polymerization [3, 4]. While controlled polymerizations are currently being converted into industrial-scale processes, the recent progress of polyolefin synthesis has revolutionized polyolefin production and is enabling the controlled formation of a wide variety of polyolefin materials [5]. Modern, highly active, selective catalysts have greatly simplified olefin polymerization and eliminated the need for extensive purification related to removal of catalyst residues, wax-like oligomers, and steroirregular byproducts. They produce more than 10 tons of polyolefin per gram of titanium – catalyst residues can be left in the polymer. Comonomer incorporation and the design of multimodal molar mass distribution have led to the production of linear low-density polyethylene films exhibiting an improved toughness/stiffness balance and heat-sealing capability. As a function of the catalyst architecture it is possible to control polyolefin particle morphology. In Basell’s Spheripol process, pellet-sized polypropylene is obtained in liquid propylene using spherical catalysts as templates. The average size of a single polyethylene reactor increased from 50 000 tons per year in the 1980s to more than 600 000 tons per year today. The dramatic increase in polyolefin production is accompanied by reduced margins. Strategic global alliances were formed to cut down costs and make polyolefin-large scale production more effective. The better insight and mechanistic understanding of catalytic olefin polymerization has led to novel generations of single-site catalysts (see Fig. 13.2). Whereas the first catalyst generation consisted of a large number of catalytically active transition metal compounds exhibiting different reactivities and selectivities, the new generation of catalysts contains only one type of catalytically active sites which can be fine-tuned with respect to molar mass control, end-group for-

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Fig. 13.2 Polymerization of olefins using multi- and single-site catalysts such as metallocenes (left).

mation, and comonomer incorporation. Typical examples are metallocene-based single-site catalysts [6] and new “beyond-metallocene” catalysts containing rigid ligand frameworks [7]. Conventional catalysts gave a complex mixture of copolymers with predominant incorporation of comonomer in a wax-like, low molecular weight fraction. The novel single-site catalysts can be tuned to give unprecedented control of the molecular architectures, e.g., for random comonomer incorporation over the entire composition range including low-reactivity comonomers such as 1-octene, styrene, and cycloolefins. Also included are thermoplastic elastomers such as ethane–1-octene copolymers and rubbers (Engage, from Dow). Block copolymers can also be formed using the specialized beyond-metallocene catalyst generations [7]. Variation of the polypropylene stereochemistry can produce stiff isotactic PP and also flexible stereoblock elastomeric PP. A wide range of oligomers has become available. Clariant’s metallocene-based waxes (Licocene) are used as dispersing agents and compatibilizers in various formulations. Some of the emerging modern catalyst systems exhibit improved tolerance to polar comonomers which cause severe catalyst poisoning of most state-of-the-art catalysts. Late transition metal catalysts polymerize ethylene in water and produce polyethylene emulsions in catalytic nano-suspension polymerization [8]. Further breakthroughs are expected to open new routes for tailor-made adhesive components based upon polyolefin feedstocks. 13.2.2 Functional Polymers from the Life Sciences

In nature the enzymes are able to convert cheap feedstocks such as sugars and amino acids into a large variety of functional and structural polymers with very high complexity. For many years, biopolymers such as starch, dextrose, cellulose, shellac, casein plastics, and proteins were used as polymers from renewable resources to formulate adhesives [9]. The life sciences effort is promoting the development of new processes based upon biotechnological routes to carbo-


hydrates, polyesters, and phenolic resins [10]. Biotechnology routes have been established to produce thermoplastic polyhydroxybutyrate (Biopol) and polylactide (NatureWorks from Cargill), some of which can also be obtained from petrochemical feedstocks using conventional catalysts together with propylene oxide and carbon monoxide [11]. Biotechnology is being exploited in lignocellulose processing [12] and manufacture of protein-based marine adhesives [13]. Without any doubt the remarkable progress in programmed biosynthesis will lead to the production of functional polymers useful in adhesive formulations. This “green polymer chemistry” offers the unique potential to produce complex functional polymers in one step. 13.2.3 Reactive Extrusion and Isocyanate-Free Polyurethane Chemistry

In addition to catalysis and controlled polymerization, the traditional advancement reactions will continue to play an important role. Novel chain extenders are becoming available which facilitate the controlled formation of molecular architectures with tailored property profiles. A prominent example of tailor-made polymers is polyurethane, which requires toxic isocyanates for the formation of block copolymers. As a result of poor selectivities, a large excess of diisocyanates must be added in order to make oligomers with reactive end groups. Since excess isocyanates as well as some of their reaction products are toxic, special purification is needed. In a recent approach, the new isocyanate-free nontoxic chain extender carbonylbiscaprolactam (CBC, Allinco; DSM; see Fig. 13.3) became available to produce caprolactam-blocked isocyanates, polyurethanes, and polyureas, as well as polyester urethanes and polyester ureas in melt phase reactions [13–16]. The side product caprolactam is readily removed by vacuum stripping. The conversion of primary amines and hydroxyl groups is highly selective. At 100 8C bulk conversion of primary amines occurs quantitatively in the presence of secondary amines. Highly branched, blocked isocyanate resins are available from polyamines and polyols (Scheme 13.1). Since nontoxic components are used, biocompatible polymers are available for biomedical applications. A large variety of reactive oligomers and highly functionalized intermediates for

Fig. 13.3 Carbonylbiscaprolactam chain extender (Allinco, from DSM).

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Scheme 13.1 Polyol conversion to blocked isocyanates without using either phosgene or isocyanate. CBC = carbonylbiscaprolactam.

adhesive and coating applications as well as for in-situ primers are available via this versatile melt phase coupling route.

13.3 Nanosystems 13.3.1 The Nano Challenge

Since the 1990s nanotechnology and nanomaterials have been attracting considerable attention in academia and industry as well as in investment banking [17]. The technical term “nanotechnology” was listed in the “Science Finder” database for the first time in 1978. It should be noted, however, that the technical term “colloid” was first introduced as early as 1900 to describe special silver nanoparticles used in photographic applications. In rubber processing, nanometer-scale fillers have a very long tradition. Various nanofillers were applied in coating and adhesive formulations. Pyrogenic silicas of 5 nm diameter are well-known thixotropic additives which control rheology via their shear-sensitive self-assembly. Nanometer-scale phases are obtained either by nanophase separation processes or by dispersion of nanometer-scale particles or nanomolecules, respectively. The formation and assembly of nanoparticles and nanomolecules within a polymer matrix is displayed in Fig. 13.4. When microparticles are substituted by the equivalent volume fraction of nanoparticles, the average particle number in-

13.3 Nanosystems

Fig. 13.4 Nanoparticle assembly and gradient materials.

Fig. 13.5 Assembly of nanoparticles and nanomolecules (“tectons”).

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creases by nine orders of magnitude. As a consequence, the entire polymer accumulates at the nanoparticle surface, even at low weight fractions of nanoparticles of just a few weight percent. This represents an attractive route to convert bulk polymers into interfacial polymers with new property profiles without changing the monomer compositions. Moreover, the assembly of nanoparticles can produce skeleton-like superstructures or gradient materials. This gradient formation can lead to in-situ improvement of surface and interface adhesion, improved fracture toughness via multiple plastic deformation, enhanced barrier performance by increasing the diffusion pathway, and increased dimensional stability. The nanometer dimension is bridging the gap between materials and molecules! This is of special interest for the formation of organic/inorganic hybrid materials. In Fig. 13.5 nanometer-scale building blocks (“tectons”) are displayed which can be dispersed and assembled within the polymer matrix. 13.3.2 Nanophase Separation

It is quite well known that the formation of nanophases plays an important role in adhesive technology although this fact was ignored for many years due to the difficulties relating to the imaging of such small structures. Nanometer-scale interdiffusion layers account for polymer/polymer adhesion. This is illustrated in Fig. 13.6 for the sandwiched films of the thermoplastic elastomer SEBS and isotactic polypropylene, annealed at 160 8C for several hours. The interdiffusion layer is approximately 100 nm wide. This interfacial nanodesign is the key to improved adhesion of polypropylene materials. The formation of rubber nanophases can improve the toughening of rubbermodified adhesives. For many years the role of nanophases was ignored because they were not detected by conventional microscopic imaging. In fact, most of the effort in the development of liquid nitrile rubber was put into microphase separation. The impact of nanostructure formation was discovered during the early 1980s. When epoxy-terminated nitrile rubber was blended together with phenolic oligo(ether urethane)s derived from oligo(propylene glycol)s, unusual rubber blend toughened epoxy structural adhesives were obtained [18]. Only when the

Fig. 13.6 The transmission electron microscopic (TEM) image of an SEBS/iPP interface.

13.3 Nanosystems Fig. 13.7 The TEM image of the nanophase-separated structural epoxy adhesive containing nitrile rubber/phenolic polyurethane liquid rubber blends (cure with dicyandiamide at 180 8C).

compatibility was matched, were nitrile rubber nanophases formed together with nanophase-separated polyurethane rubber phases (see Fig. 13.7). This approach to epoxy toughening afforded unusual synergisms of simultaneously improved toughness, stiffness, strength, and adhesion. As is apparent from Fig. 13.8, the T-peel resistance of bonded steel increased with increasing strain rate until the steel substrate was destroyed. In conventional toughened epoxy adhesives, containing exclusively the microphase-separated nitrile rubber, the T-peel resistance declined with increasing strain rates. The fatigue resistance of the liquid rubber blend system was also improved substantially. This adhesive generation was devel-

Fig. 13.8 The T-peel resistance as a function of the crosshead speed for nitrile rubber/phenolic polyurethane liquid rubber blend toughened epoxy (upper curve) and conventional nitrile rubber-toughened epoxy structural adhesives (lower curve).

Fig. 13.9 Box beam bonding using rubber blend toughened epoxy adhesives.

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oped at Ciba and introduced into commercial structural bonding by Gurit-Essex and Dow. Such nanophase-separated toughened epoxy adhesives can be employed as high-performance adhesives. Fig. 13.9 shows an epoxy-bonded box beam which dissipated the impact energy encountered during a car crash. 13.3.3 Nanomolecules as Molecular Nanoparticles

Most linear polymers change conformation as a function of solvent polarity and shear forces. In contrast, highly branched, tree-like polymers, also referred to as dendrimers or hyperbranched polymers (Fig. 13.10), have a spherical shape which is independent of shear, pH, and solvents. Although nature has been exploiting hyperbranched polymers such as amylopectin for millions of years, the consequent design of dendritic or hyperbranched polymers was started only in the 1980s. Interesting features of such polymers are: (a) low viscosity even at high molecular weight due to their spherical shape; (b) high solubility and molecular dispersion in various thermosets; (c) core/shell topology with highly functionalized surfaces; (d) polarity and surface design via end-group variation; (e) controlled self-assembly in the bulk and at surfaces. Several dendritic and hyperbranched polymers are available on a commercial, semi-commercial or pilot scale: · dendritic polyamidoamine from Tomalia [19, 20] · dendritic polypropyleneimine (Astramol from DSM) [21] · hyperbranched polyamidoamine (Dendrepox from Epoxy Ltd.) · hyperbranched polyesteramide (Hybrane from DSM) [26–29] · hyperbranched polyester based upon dimethylolpropionic acid (Boltorn from Perstorp Co.) [22–24]

Fig. 13.10 Dendrimers (left) and hyperbranched polymers (right).

13.3 Nanosystems

· hyperbranched polyethyleneimine (BASF AG) · hyperbranched polyols and polyurethanes (BASF AG) [25] · hyperbranched polyglycerol from gylcidol (Hyperpolymers GmbH) [21]. While true perfect dendrimers such as Tomalia’s PAMAM [19, 20] require multistep syntheses involving large excesses of reagent, most of the less-perfect hyperbranched polymers are readily available in one-step, one-pot syntheses. For example, AB2+x-type monomers are used in polyaddition or polycondensation processes. Living multibranching anionic ring-opening polymerization of glycidol afforded high molecular weight hyperbranched polyglycidol in very high yield and high purity [21]. The molecular topology and the corresponding AFM image are shown in Fig. 13.11. Nanometer-scale molecules of hyperbranched polymers and dendrimers can be dissolved and assembled in various polymer matrices as well as on polymer surfaces and polymer/polymer, polymer/filler and polymer/fiber interfaces. They provide attractive properties: · improved adhesion · high bond strength · effective toughening · a low-stress matrix · variation of cure chemistry · interpenetrating network formation. The syntheses, properties and applications of hyperbranched polyesters have been investigated in great detail [22]. The addition of hyperbranched polyesters was reported to double the interlaminar fracture resistance of epoxy-based composites and to reduce the internal stress level by as much as 80% in the pres-

Fig. 13.11 Hyperbranched polyglycidol as a molecular nanoparticle, with its core/shell topology.

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ence of only 10 wt.% modifier. Due to the nature of the molecules, filtration during fiber impregnation can be prevented in composite formations [23]. Hyperbranched polyesters can also be employed as compatibilizers to improve the dispersion of polypropylene/polyamide blends [24]. The potential of hyperbranched polymers as highly functional, in-situ, interfacial and surface adhesion promoters was confirmed by many other groups. Bruchmann at BASF introduced new families of hyperbranched polyols and polyurethanes as very effective built-in adhesion promoters, useful in coatings and film printing [25]. Similar surface anchoring mechanisms were proposed to explain the very high effectiveness of hyperbranched polyesters, polyesteramides, and polyimines as dispersing agents of pigments [26]. Hyperbranched polymers and dendrimers are effective nanocarriers and hosts for dyes, pigments, and other nanoparticles which are readily dispersed [27]. At DSM new families of high-performance additives based upon dendritic Astramol and hyperbranched Hybrane polymers were introduced [28, 29]. 13.3.4 POSS and Nanocomposites

When the size of solid inorganic materials is reduced to approach the nanometer range, there exists an interesting dualism between molecules and particles. In fact, it no longer makes sense to distinguish between molecules and particles. For example, the sol–gel reaction of functionalized trisalkoxysilanes can be tuned to produce polysilsesquioxanes which exhibit nanometer dimensions and can be dispersed and assembled within the matrix of epoxy and acrylic resins [30]. A wide variety of polyhedral oligomeric silsesquioxanes (POSS) are offered by Hybrid Plastics to close the gap between inorganic solid and hybrid nanomolecules. Today a large variety of inorganic nanoparticles are available. Two approaches are being exploited to produce nanocomposites with improved property profiles: (a) dispersion and (b) intercalation. In the dispersion route, nanoparticles are rendered organophilic by means of surface modification. The two classes of epoxy nanocomposites are displayed in Fig. 13.12. Spherical nanosilicate dispersions are obtained from sodium silicate nanoparticles which are rendered organophilic by means of surface sodium-cation exchange (acid activation with ion-exchange resins) combined with organophilic silane surface modification. Dispersions of uniform spherical particles with sizes varying between 20 and 500 nm are readily available. In the second approach, layered silicates are intercalated. For example, clay mineral contains anionic aluminosilicate layers which are rendered organophilic by means of cation exchange of the sodium interlayer cations for organophilic alkylammonium cations. When the compatibilities of a polymer and an organophilic layered silicate are carefully matched, diffusion of the polymer between the layers causes intercalation and ultimately exfoliation, as shown in Fig. 13.12 for the dispersion of the individual nanoplatelets. According to Kinloch, the addition of 1–8 wt.% of nanosilica particles with an average diameter of around 20 nm to a typical rubber-toughened adhesive, based

13.4 Conclusion

Fig. 13.12 Polymer nanocomposites containing spherical silica nanoparticles in vinyl ester (left) and exfoliated organoclay in polypropylene (right).

upon a two-part epoxy formulation, leads to very significant increases in the toughness of the adhesive and also to increases in the glass transition temperature and the single-lap shear strength. The significant improvements of the mechanical properties are accompanied by improved thermal stability [31]. In the case of montmorillonite and fluorohectorite, which are rendered organophilic by means of cation exchange of sodium cations for various alkylammonium compounds, significant improvements in stiffness and toughness can be obtained, provided that the alkyl chain length of the alkylammonium cation-exchange reagent has more than 12 methylene units and that adequate compatibility between surface modification and resins is achieved [32, 33]. The alkylamine adduct formation appears to facilitate diffusion of epoxy resins between the silicate layers. Several nanocomposite families have been reported to offer attractive potential for improving the performance of adhesives [34]. The combination of organophilic layered silicates and compatibilized liquid rubbers can be exploited to improve toughness and stiffness simultaneously [35].

13.4 Conclusion

The recent progress in polymer synthesis and nanotechnology is stimulating the development of adhesives with improved performance. A few percent of nanometer-scale additives can be added to formulations in order to achieve significant changes in property profiles. Molecular nanoparticles and nanometerscale, highly branched polymers can be dispersed in order to facilitate energy dissipation at the crack tip by means of multiple plastic deformation. The selfassembly of nanoparticles forms skeleton-like superstructures which account for

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improved dimensional and thermal stabilities and barrier performance. The self-assembly of functionalized nanoparticles and nanomolecules at nano-reinforced surfaces and interfaces (nano-armoured interfaces) provides very effective adhesion via covalent or ionic groups. The precise control of surface functionality and the attractive particle/molecule dualism can be exploited to design highly functional resins without a buildup of viscosity, and to prepare novel organic/inorganic nanohybrid materials combining the prospective benefits of polymer and ceramic materials, e.g., an improved toughness/stiffness balance, high adhesion, an effective moisture barrier, oil adsorption and retention, corrosion resistance, reduced thermal expansion, and halogen-free flame retardancy. Nanoadditives are the key components of high-performance adhesives and will promote future progress in adhesive technology.

Acknowledgments

The author gratefully acknowledges financial support by the Deutsche Forschungsgemeinschaft (DFG). The TEM image of SEBS/iPP was prepared by Ralf Thomann and Professor Joerg Kressler at FMF.

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Mülhaupt, Macromolecules 1999, 32(13), 4240. a) A. Hult, M. Johansson, E. Malmstrom, Adv. Polym. Sci. 1999, 143 (Branched Polymers II), 1–34; b) S. Nummelin, M. Skrifvars, K. Rissanen, Topics Curr. Chem. 2000, 210 (Dendrimers II), 1–67. R. Mezzenga, L. Boogh, J.-A. E. Manson, Composites Sci. Technol. 2001, 61(5), 787. G. Jannerfeldt, L. Boogh, J.-A. E. Manson, Polymer 2000, 41(21), 7627. B. Bruchmann, GIT Labor-Fachzeitschrift 2002, 46(12), 1346. F. O. H. Pirrung, E. M. Loen, A. Noordam, Macromolecular Symposia 2002, 187 (Quo Vadis – Coatings?), 683–693. P.E. Froehling, Dyes and Pigments 2001, 48(3), 187. P. Roehling, P. J. Brackman, Macromolecular Symposia 2000, 151 (Polymers in Dispersed Media), 581–589.

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Polym. Polym. Compos. 1999, 7(8), 581– 587. E. Müh, H. Frey, J. E. Klee, R. Mülhaupt, Adv. Functional Mater. 2001, 11(6), 425. A. J. Kinloch, J. H. Lee, A. C. Taylor, S. Sprenger, C. Eger, D. Ega, J. Adhes. 2003, 79(8–9), 867. C. Zilg, R. Thomann, J. Finter, R. Mülhaupt, Macromol. Mater. Eng. 2000, 280/ 281, 41–46. X. Kornmann, R. Thomann, R. Mülhaupt, J. Finter, L. Berglund, J. Appl. Polym. Sci. 2002, 86(10), 2643–2652. A. Hartwig, D. Putz, M. Sebald, Adhäsion – Kleben & Dichten, 2003, 47(1–2), 16–18, 20. J. Fröhlich, R. Thomann, R. Mülhaupt, Macromolecules 2003, 36(19), 7205–7211.

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14 Influence of Proton Donors on the Cationic Polymerization of Epoxides A. Hartwig, K. Koschek, and A. Lühring

Abstract

The cationic polymerization of epoxy resins becomes more and more important in adhesive bonding technology. Most often a cationic photopolymerization is applied, but occasionally thermally induced cationic polymerization of epoxides is also used. The polymerization is strongly influenced by water, namely the humidity of the surrounding atmosphere in which the adhesives are cured. Traces of water or other proton donors are required as co-catalyst for the initiation of the polymerization. On the other side the cation as well as the anion of the latent initiators is hydrolyzed by the long term influence of moisture. But water does not only influence the initiation reaction of the polymerization but also the chain growth reaction. Depending on the kind of epoxide the polymerisation is strongly accelerated by the water or strongly retarded. These extreme differences are explained by differences in the mechanism of the chain transfer reaction carried out by the water. In any case the water reduces the mechanical strength of the cured epoxides by the formation of unbound chain-ends in the polymer network and therefore lower cross-link density. Also alcohols lead to chain transfer reactions during polymerization of the epoxides and consequently they have an influence on the polymerization kinetics as well as on the mechanical properties of the epoxides Different polymeric and oligomeric alcohols (polyols) are applied to adjust the mechanical properties of cationically curing adhesives. It could for example be shown that the addition of polytetrahydrofuran leads to polymers with rubber elastic behaviour. But also the polymerization kinetics is strongly influenced by polyols. In the case of polyethers based on 1,2-diols the protons required to progress the polymerization can be fixed within five-membered rings. The polymerization rate is decelerated in the presence of such polyols, therefore. The formed structures have some similarity with crownethers complexing a proton. And it is therefore not surprising that crown-ethers inhibit the cationic polymerisation very efficiently. On the other hand no retardation of the reaction rate is observed if 1,4-diol based polyethers are used to modify the properties of the materials. Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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14 Influence of Proton Donors on the Cationic Polymerization of Epoxides

14.1 Introduction

The cationic photopolymerization of epoxides is widely used for printing inks, adhesives, and coatings. Appropriate initiators are diaryliodonium and triarylsulfonium salts of so-called superacids. Common counterions are PF–6 and particularly SbF–6 [1, 2]. An interesting, industrially applied anion is [B(C6F5)4]–, which gives the salts a better solubility in hydrophobic media – especially epoxy silicones – and during hydrolysis there is no formation of fluoride ions. This is of interest for the application of the resulting reactive mixtures in electronics or in the bonding of glass, as well as for basic investigations and research, because the formation of fluoride ions from purely inorganic anions is thought to be the termination reaction in the living cationic polymerization of epoxides. During the last few years latent thermal initiators for the cationic polymerization of epoxides were also developed [2]. The organic cations are largely ammonium, sulfonium, and iodonium ions. The polymerization of the epoxide is initiated by the superacid or carbocation formed by the photochemically or thermally induced decomposition of the initiator. The basic reactions for the photochemical decomposition of diaryliodonium [3] and triarylsulfonium salts [4], as well as the initiation, chain growth, and chain transfer of the epoxide polymerization [5], are described in the literature. The effect of water and alcohols on the polymerization of epoxides with latent initiators has received little attention, although it is well known that hydrogen donors strongly influence the cationic polymerization of any monomer. If the reaction is not directly initiated by the carbocations formed by the decomposition of the initiator, a proton donor is required as co-catalyst for the formation of the initiating superacid [3, 4]. As long as the monomer is not dried extensively and the reaction is not carried out in a dry atmosphere, the amount of water naturally present in the monomer is high enough to work as a proton donor for the initiation. Furthermore the presence of water or alcohol leads to chain-transfer reactions. The alcohol attacks the positively charged end of the growing polymer chain, forming an ether bond, and in the case of water as reaction partner a hydroxyl group is formed. The released proton initiates the growth of the next polymer chain. This is also called “polymerization by the activated monomer mechanism” and was examined with conventional initiators by Penczek et al. [6]. Analogous work for the photochemically induced polymerization of epoxides was published by Yagci and Schnabel [7], who showed that a kind of copolymerization takes place between the diepoxide and the alcohol via a chain-transfer reaction. If a diepoxide is polymerized with a diol or triol, the system becomes crosslinked, whereas the reaction with a monoalcohol leads to dead ends in the polymer network [8]. Crivello et al. [9] found that the rate of the photochemically induced polymerization of 3,4-epoxycyclohexylmethyl-3',4'-epoxycyclohexane carboxylate increases if ethylene glycol or 1,4-butanediol is added. At the same time the conversion of epoxy groups increases within the first 3 min of UV irra-

14.2 Initiators for the Cationic Polymerization of Epoxides

diation. The curing rate was examined semiquantitatively using the irradiation time required to receive a tack-free coating. Corresponding measurements for the cationic photopolymerization of vinyl ethers showed that for an amount of up to 5 wt.%, alcohols did not influence the polymerization rate, but they increased the final conversion of double bonds [10]. Since the network density is changed by the reaction between epoxide and alcohol or water, the mechanical properties of the resulting polymer are also influenced. This can be used, for example, in the flexibilization of dental materials with poly(1,4-butanediol) [11] or coatings with polyester polyols [12]. It is not only alcohols that influence the polymerization kinetics and the properties of the polymer, but also carboxylic acids. By the addition of a polymer with carboxylic acid groups instead of the polyol, a polyester is formed as a reaction product and not a polyether. This was examined in detail by Wu and Soucek [13]. After this introductory review on the cationic polymerization of epoxides by latent initiators an overview of our work in recent years on the effect of moisture and different polyether diols on the photochemically and thermally induced polymerization of different epoxides will be given. The polymerization behavior, as well as the final properties of the cured polymers, is covered.

14.2 Initiators for the Cationic Polymerization of Epoxides

There are many initiators available for the radical photopolymerization of compounds such as acrylates. This is not the case for photoinitiators for the cationic polymerization of epoxides. However, the few available sulfonium and iodonium salts, which are used in the following investigations, have several disadvantages. The sulfonium salts are commercially available as 50% propylene carbonate solutions and, during the polymerization of epoxides, the solvent is built into the polymer network via a ring-opening reaction. From an application point of view, the use of such a compound as solvent is quite practical, since it affords a quick homogeneous distribution of the initiator in the monomer. For the basic investigations, especially those into the influence different components have on the polymerization kinetics, the propylene carbonate would cause problems. The iodonium salts described in the literature [1, 2, 5] or those commercially available have the disadvantage that in practice they only absorb 250–270 nm radiation, and would therefore not be activated by normal UV radiation. For industrial applications this disadvantage is removed by the addition of photosensitizers, especially anthracene, perylene, benzophenone, 2,2-dimethoxy-1,2-diphenylethanone, and 2-isopropyl-9H-thioxanthen-9-one [14]. The effectiveness of an intermolecular sensitization and therefore the progress of the polymerization reaction being investigated should become lower with increasing viscosity. This would have an unknown influence on the polymerization reactions under investigation. It is also expected that with similar effectiveness a larger amount of initiator is needed for an intermolecular than for an intramolecular reaction. At-

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Scheme 14.1 Synthesis of (9-oxo-9H-fluoren-2-yl)phenyliodonium hexafluoroantimonate.

tempts were made to develop an iodonium salt which shows sufficient absorption in the emission region of a mercury vapor lamp (300–400 nm). With this aim fluorenone was converted to an iodonium salt according to Scheme 14.1. Using real-time IR spectroscopy it could be shown that this initiator induced a faster epoxide polymerization than a commercial iodonium salt alone or with different sensitizers. A sensitization of the commercial iodonium salt phenyl [p(2-hydroxydecanoxy)phenyl]iodonium hexafluoroantimonate with fluorenone could not be observed [15]. One of the focal points of this work was the influence of traces of moisture on the cationic polymerization of epoxides. Using the previously described iodonium salt as an example, it was shown that water not only played an important role as co-catalyst in the photochemical formation of the superacid, but it also caused the hydrolysis of the iodonium salt via a nucleophilic attack at the I+ center. This reaction was mainly investigated by X-ray photoelectron spectroscopy (XPS). Interestingly, this hydrolysis only occurs with the hexafluoroantimonate salt and not with the corresponding chloride. Here it should be mentioned that the chloride salt resembles an interhalogen compound and the hexafluoroantimonate salt forms a real ion pair. Hydrolysis of the hexafluoroantimonate anion of the iodonium salt was not observed, but slow hydrolysis was observed for sodium hexafluoroantimonate. The hydrolysis is so slow overall that the influence of this effect can be excluded from the discussions into the influences of water and alcohol on the cationic polymerization of epoxides. Due to the high effectiveness of the iodonium salt developed, it is used preferentially in the following investigations as a photoinitiator for the polymerization of epoxides. In some cases a commercial sulfonium salt is also used, in spite of the known disadvantages. In other cases the polymerization is initiated thermally. A requirement for a thermally induced polymerization with a latent initiator is that at room temperature the reaction mixture of monomer and initiator remains unchanged for as long as possible, but at moderate temperatures the polymerization should be fast. For the radical polymerization of, for example, acry-

14.3 Influence of Moisture on the Polymerization Kinetics

lates or styrene, this criterion is not achieved satisfactorily. For the thermally induced polymerization of epoxides there are, however, several compounds described in the literature with which these requirements can be achieved. Examples are sulfonium, ammonium, or pyridinium salts with non-nucleophilic anions [2]. With such compounds as initiator, typical reaction times at 80 8C are around 30 min. Within this time the reactive mixture polymerizes to a network polymer and despite this high reactivity the mixtures are stable for several weeks at room temperature. Usually such compounds are organic ammonium or sulfonium salts of the superacids described. Initiators with the highest reactivity are achieved when hexafluoroantimonate is used as anion. For investigations into the thermal polymerization of epoxides, benzyltetrahydrothiophenium hexafluoroantimonate [16] was chosen as the initiator; not only does it display a good balance between reactivity and stability of the reactive mixture, but it also has good solubility in the chosen epoxides and it is prepared via a simple synthesis. Another method for the thermally induced cationic polymerization of epoxides is the redox activation of the iodonium or sulfonium salts commonly used as photoinitiators. This can be achieved with, for example, copper(II) salts or ascorbic acid derivatives [17]. It could be shown that even the commercially available salt (tolylcumyl)iodonium tetrakis(pentafluorophenyl)borate with [B(C6F5)4]– as the anion could be activated using ascorbic acid- 6-hexadecanoate [18]. The onset temperature of the polymerization was determined using differential scanning calorimetry (DSC) and lies around 110 8C. This initiator system is also used in the following investigations.

14.3 Influence of Moisture on the Polymerization Kinetics

Adhesives, varnishes, and paints which polymerize via the cationic photopolymerization of epoxides are often advertised with the argument that in contrast to the radical photopolymerization of acrylates, the polymerization reaction is not inhibited by oxygen from the air. This is partly correct, but it can be influenced by air humidity and in practice the polymerization rate can increase or decrease with increasing humidity. Whether the reaction in a specific case is accelerated or decelerated depends on numerous details of the reaction mechanisms. Firstly it concerns processes during initiation of the polymerization, to which the previously discussed hydrolysis of the iodonium salts belongs, as well as the proton donor effect during formation of the initiating hexafluoroantimonic acid from decomposition of the initiator. It can be assumed that the necessary amount of water is always present in epoxy resins which are handled in air or not specifically dried. It could be shown by real-time infrared spectroscopic (RTIR) examination of the photopolymerization of monomer layers 12 lm thick that the photochemical polymerization of the cycloaliphatic epoxy resin 3,4epoxycyclohexylmethyl-3',4'-epoxycyclohexanecarboxylate does not occur in dry air if the epoxy resin does not contain alcohols [19]. This can mean either that

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initiation does not occur via the carbocation formed, or that chain growth is not possible. In the second case it would mean that the published polymerization mechanism of a chain-growth reaction for this monomer is incorrect and consumption of the monomer occurs instead via a chain-transfer reaction, which will be discussed below. If initiation occurs with hexafluoroantimonic acid and not a carbocation, then hydroxyl groups form at the chain ends. This is a requirement for the subsequent chain-transfer reactions resulting in monomer consumption. A definite distinction between the impact of these influences is, however, not possible. There is an indication in the literature that polymers form corresponding groups at the chain ends when they react with the postulated carbocations [20]. It was established in addition that layers of the monomer–initiator mixtures which were irradiated in a dry environment did not polymerize. After a wet atmosphere was introduced, polymerization occurred. This suggests that the carbocation formed could initiate the reaction, but monomer consumption could only occur through a chain-transfer reaction. These findings correspond quite well with the results of Decker [21]. It was shown with RTIR [19] that for epoxides with exocyclic epoxide groups the rate of polymerization is not influenced by air humidity (vinyl epoxides), or decreases with increasing air humidity (glycidyl ethers). For epoxides with endocyclic epoxide groups an increase in the rate of polymerization is observed with increasing air humidity. The reason for the different influences of air humidity on the polymerization rate for different epoxides is differences in the chaintransfer reaction carried out by water. In the case of glycidyl ether both water protons are fixed in a five-membered ring at the positively charged end of the growing chain (Scheme 14.2). Through this fixation and the delocalization of the positive charge, the release of a proton, a requirement for the start of a new chain reaction, is hindered. This can lead to the observed retardation of the polymerization reaction with increasing humidity. During the attack of water on the growing chain end of an epoxide polymer that is not a glycidyl ether, e.g., an epoxide with endocyclic epoxide groups, only one of the attacking water protons is fixed in a five-membered ring (Fig. 14.1). The second proton can be released without having to overcome an energy barrier. In this case the chain-transfer reaction is not retarded. This, however, does not explain the increasing polymerization rate with increasing air humidity, since water should only slightly retard or should not influence the polymerization rate. Molecular modeling was applied to explain this. In Fig. 14.2 it can be seen that the front side of the protonated epoxide group is perfectly shielded by

Scheme 14.2 Chain-transfer reaction by water for the polymerization of a glycidyl ether.

14.3 Influence of Moisture on the Polymerization Kinetics Fig. 14.1 Optimized conformation of a polymerizable endocyclic epoxide after reaction of the chain end with water (the growing chain end is assumed to be a methyl group). Bond lengths are given in Ångstroms.

the SbF–6 ion. It is not possible for the attack of another epoxide to occur and therefore a chain reaction cannot begin. Small molecules, e.g., water, can attack from below or from the rear, so that epoxide consumption occurs. The resulting protonated alcohol is sterically less hindered than the protonated epoxide (each with counterion), so that an epoxide ring can be attacked by this compound. This partial step of the chain-transfer reaction leads to an acceleration of the polymerization with increasing humidity. The polymerization of epoxidized vinyl groups is practically uninfluenced by water. This is because the epoxide is accessible from at least one side, and can therefore be polymerized more quickly. Due to this high reaction rate a reaction with water cannot occur. In the case of a reaction with water, the proton is released quickly and this does not lead to a retardation of the chain-transfer reaction [19]. The presence of water during polymerization does not influence the polymerization kinetics alone, but also the structure and mechanical properties of the polymer network formed. During polymerization water is built into the polymer network in considerable amounts. By examination of the polymerization with RTIR it could be shown that hydroxyl groups were formed; in the case of cycloaliphatic epoxy resins there is even a linear relationship between turnover and the formation of the hydroxyl groups [19]. Through hydroxyl group formation, the ends of the chains in the polymer network remain unattached. During poly-

Fig. 14.2 Shielding of a protonated endocyclic epoxide group by a hexafluoroantimonate ion. Bond lengths are given in Ångstroms.

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merization the network density decreases with increasing air humidity. The result is that the polymer formed becomes increasingly soft with increasing humidity. This could be shown by measuring the hardness of small drops of a cationically polymerized bisphenol A diglycidyl ether [22]. Since this corresponds to a living polymerization, the resulting hardness does not depend on the irradiation time as long as there is sufficient waiting time between irradiation and measurement.

14.4 Modification of the Polymerization Behavior by the Addition of Alcohols

Analogously to water, alcohols (used in the form of polyols) also cause a chaintransfer reaction during the cationic polymerization of epoxides. Polyols (polyvalent alcohols with an average molecular weight from several hundred to several thousand grams/mole) are often utilized in technical formulations. These serve to reduce the network density of the polymerized epoxide and therefore make the material less brittle. At the same time the reactivity also is influenced and the susceptibility of the polymerization rate and polymer properties to the influence of air humidity is reduced. This is important, as the influence of air humidity is difficult to reproduce in technical applications. Extensive investigations with different polyether polyols and the cycloaliphatic epoxy resin 3,4-epoxycyclohexylmethyl-3',4'-epoxycyclohexane carboxylate show that either they do not influence or they decrease the rate of polymerization [23]. Due to the higher molecular mobility during polymerization, and therefore vitrification at a higher turnover for a given polymerization temperature, as well as the action as a chain-transfer agent, a greater conversion within the investigated time period is found for the cases where alcohols were added compared to those with just the pure epoxide. The 1,2-diol-based polyethers poly(ethylene glycol) and poly(propylene glycol) lead to a reduction in the polymerization rate, whereas the 1,4-diol-based polyether polytetrahydrofuran has practically no influence on the rate of polymerization; only the conversion is increased. Similarly to the investigations into the influence of air humidity, the corresponding ones into polymerization kinetics of the epoxide layers on gold substrates were carried out with realtime IR spectroscopy. The difference in influences of the 1,2-diol-based polyethers compared with the 1,4-diol-based polyether can again be explained by the fixation of protons in a five-membered ring. In both types of polyether the chain-transfer reaction occurs at the hydroxyl end group of the diol. During reactions with 1,2-diol-based polyethers the reactive proton can be immobilized through several hydrogen bonds and its availability to start the next chain reaction is reduced (Scheme 14.3). This leads to the observed reduction in reaction rate. This type of additional hydrogen bond between the oxygen of the polyether and a proton cannot be formed in 1,4-diol-based polyethers. Therefore, the presence of these polyethers does not lead to a decrease in reaction rate.

14.4 Modification of the Polymerization Behavior by the Addition of Alcohols

Scheme 14.3 Proton fixation through hydrogen bonds and proton release as the last step in a chain-transfer reaction with a 1,2-diol-based polyether.

Apart from the suppression of the chain-transfer reaction, the retardation of the reaction for 1,2-diol-based polyethers can also be explained by the interaction of protons with the polyether chain. A proton can be bound to the polyether chain with a corresponding conformation via several hydrogen bonds (Fig. 14.3, left). The resulting structure has similarities to a crown ether with a bound proton (Fig. 14.3, right). This leads to a reduction in the concentration of the available superacid, which is needed to initiate the polymerization. Beside the photochemically induced polymerization described in Section 14.2, the reaction was also induced using benzyltetrahydrothiophenium hexafluoroantimonate as initiator. It could be shown with DSC investigations that the onset temperature of the reaction and the temperature of the reaction maximum increased with increasing chain length from ethylene glycol to poly(ethylene glycol) 600, and that even the addition of a small amount of 12-crown-4 to a reaction mixture of cycloaliphatic epoxy resin and polytetrahydrofuran greatly increases the onset temperature. The temperature for the reaction maximum is hardly influenced. If the crown ether is added to a mixture which contains poly(ethylene glycol) instead of polytetrahydrofuran, then the reaction is not influenced. The investigations show that the interaction of a proton with 1,2-diolbased polyether chains plays an important role in the observed retardation of the polymerization reaction, compared with samples with 1,4-diol-based polyethers. Whether the interaction of the protons with the polyether chain or the proposed fixation of the protons during the chain-transfer reaction is more important for the observed retardation of the polymerization in the presence of 1,2-diol-based polyethers could not be determined from the investigations which were carried out. Increasing the onset temperature for a reaction without increasing the temperature for the reaction maximum, by using small amounts of crown ether instead as stabilizer of the usual tertiary amines, allows the stabilization of the

Fig. 14.3 Proton fixation in a ring-like chain section from poly(ethylene glycol) and in 12-crown-4.

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cationic polymerization reaction mixture consisting of an epoxy resin, initiator, and possibly polytetrahydrofuran. Tertiary amines have the disadvantage that they can also initiate an anionic polymerization, and can therefore lead to premature gel formation of the reactive mixture. The mechanical properties of the polymer formed are influenced not only by water built into the polyether network during the chain-transfer process but also by the integration of polyols. Usually these are added to form a less brittle polymer with a lower glass transition temperature. During the polymerization of the cycloaliphatic epoxy resin 3,4-epoxycyclohexylmethyl-3',4'-epoxycyclohexane carboxylate with polytetrahydrofuran (PTHF), the resulting polymer is not only soft but also has the properties of a typical elastomer [24]. The elastomeric properties are due to a phase separation between epoxide-rich and PTHF-rich phases. In the PTHF-rich phase, interactions between chains of the crystallizable polymer are present. This corresponds to the phase separation present in urethane elastomers, which is the main cause for the specific properties of these polyurethanes. Although crystalline domains could not be identified either by X-ray diffraction or by polarization microscopy, the presence of two glass transition temperatures and a miscibility gap between the components suggests the existence of a phase separation. For mixtures containing 25 mol% PTHF, the lower Tg is in the region of 10 8C and decreases slightly with increasing molecular weight (variation range Mn 650–2900 g mol–1) of the PTHF. The miscibility gap is found at room temperature for the samples containing 25 mol% PTHF of molecular weight from about 1400 g mol–1 and higher. At the polymerization temperature all the samples are clear. It was observed that for the polymerized samples containing 25 mol% PTHF the E modulus decreased with increasing molecular weight of the PTHF from 8 to 2 MPa, whereas the elongation at break increased from 15 to 25%. The alteration of both properties shows that the network density decreases with increasing molecular weight of the PTHF. The increased elongation at break is probably due to a phase separation which becomes stronger with increasing molecular weight. This is expected not only on a molecular level but can also be observed at room temperature in the clouding, or even phase separation, of the mixture before polymerization. The thermal stability of these epoxide-based elastomers is too low for a technical application. During the polymerization elimination of tetrahydrofuran occurs, which is due to an acid-catalyzed ether cleavage. The PTHF is not only released during polymerization but also during long-term storage at higher temperatures. The amount of tetrahydrofuran which is released increases with increasing molecular weight of PTHF. It can therefore be concluded that the ether bonds between the cyclohexane ring and the PTHF are more stable than within the PTHF chains themselves. To stabilize the resulting elastomers thermally it is sufficient to add nanoparticles, namely organically modified bentonite. In this case the superacid present in the polymer is adsorbed on the surface of the nanoparticles, and is therefore not available for the decomposition of the polymer [25].

References

14.5 Conclusion

The cationic polymerization of epoxides is strongly influenced by water or alcohols. Due to their presence and the resulting chain-transfer reactions, the polymer becomes softer. This can be used to adjust the mechanical properties of the polymerized epoxides and it is even possible to prepare materials which behave like elastomers. Through the addition of proton donors, the polymerization kinetics is also strongly influenced. Depending on the resin, the presence of moisture may accelerate or decelerate the polymerization. The different behavior is explained in the different ways the proton is fixed within five-membered rings during the chain-transfer reaction. The same structures are applicable to explain the retardation of the polymerization by the presence of 1,2-diol-based polyethers or a crown ether. The knowledge gained is now being used for the development of tailor-made adhesives, potting resins, and coatings with optimized properties for specific applications.

Acknowledgments

TEM examination by H. van Eys-Schäfer and molecular modeling by B. Schneider are gratefully acknowledged.

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15 Novel Adhesion Promoters Based on Hyperbranched Polymers A. Buchman, H. Dodiuk-Kenig, T. Brand, Z. Gold, and S. Kenig

Abstract The emergence of experimental and commercial highly branched polymers (HBPs) having 2D or 3D morphologies with high peripheral functionality created new opportunities for tailoring new architectures for thermosetting adhesives. Furthermore, the HBPs found a new application as adhesion promoters on polymer, metal, and composite adherends. Our study showed that optimal adhesion properties of both shear and peel strengths were obtained when the crosslinking agents had been either partially or totally replaced by the HBPs. In addition, the insertion of HBPs into the bulk adhesive enhanced the thermal durability (Tg) due to an increase in crosslinking density and formation of modified multiphase morphologies providing high-energy absorption in the network. We also investigated the HBP PAMAM [poly(amido-amine)] system as a pre-adhesion treatment for epoxy and polyurethane adhesives and as an adhesion promoter on fibers in composite matrices. Various kinds of HBPs and dendrimers were studied on different kinds of adherend substrates, including aluminum alloy, magnesium alloy, polyetherimide (PEI), and glass/epoxy composite, and as surface treatment on fibers. The adhesion durability was characterized before and after exposure to combined heat and humidity. Experimental results have shown for various concentrations of HBPs and dendrimers that under optimum conditions the interfacial adhesion and the durability increased significantly for all the joints and composites studied. Optimum concentrations of HBPs were in the range 1–2% in alcohol, increasing the adhesion strength by 100–250% for epoxies and polyurethanes on aluminum and by 50–100% on Mg and PEI. The failure mode also changed from interfacial without treatment to mixed or cohesive following HBP treatment. HBPs improved laminate flexural strength by 60% when applied on advanced fibers such as Kevlar and Zylon, reinforcing a structural epoxy matrix.


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15 Novel Adhesion Promoters Based on Hyperbranched Polymers

15.1 Introduction

Microstructure is a key aspect for polymers in general, and for polyurethanes (PUs) in particular. The morphology of PUs is governed by the formation of hard and soft domains and their intercalation. Consequently, new microstructures can be developed using dendritic and hyperbranched (HB) polymers in PU systems. With the emergence of low-cost hyperbranched poly(amido-amine)s (PAMAMs), new avenues have been opened for enhancing the properties of epoxy and PU systems due to the modified branched network morphology. HB polymers were first postulated by Flory as polymers having a broad molecular weight (MW) distribution and non-entangled noncrystalline morphology due to their branched structure. They can be synthesized in various ways [2, 3]. The most common route is by condensation reactions carried out in either the bulk or solution. HB polymers can possess a variety of end groups (polar or nonpolar). These end groups can be reactive (e.g., acrylates, amino, or hydroxyl groups) resulting in crosslinkable HB polymers. The properties of HB polymers depend on several parameters, such as the backbone structure, the end-group chemistry, and the MW. The viscosity of HB polymers, both in solution and in the melt, is lower than that of their linear analogues, due to their branched structure. The glass transition temperature (Tg) decreases with an increasing number of end groups, and increases with increased polarity of the end groups. A higher degree of branching (DB) leads to a higher solubility and a lower melt viscosity. Tomalia [1, 8, 9] obtained, by repetitive diamine alkylation with acrylate and amidiation of the resulting ester groups, a high molecular weight (HMW) dendri-polyamide (dendrimer) [5, 6]. This method, which is a multistage process, is costly and involves a great amount of waste products. An innovative synthesis of highly branched poly(amido-amine)s was recently proposed [4]. The synthesis is based on epoxy–amine reactions (see Scheme 15.1). The branching structure is achieved by using a multifunctional core-making compound such as a tetrafunctional diamine. The process is based on two quantitative steps: diepoxy– amine addition and polyamide–epoxy addition. The second epoxy groups of the diepoxide react much more slowly with the amine groups than the first epoxy group. In this way, highly branched polymers were obtained, having an MW of 5000–15 000 g mol–1 and an amine functionality of 30–45. The secondary amine is less reactive than the primary amine, and consequently we can assume that it reacts later than the primary amine. Generally, it should be emphasized that the exact mechanism was not studied, since the reaction of epoxies and amines is well known for crosslinking and network formation. The functionality was tested by amine titration, and the molecular weight by GPC. This whole technology is based on Epox proprietary know-how, so we did not explore the very detailed mechanism, just applications of the HB products based on the patents given in the literature [7, 10–16].

15.2 Experimental

Scheme 15.1 Synthesis of hyperbranched poly(amido-amine) polymers.

The current study is aimed at investigating the effect of HB products (based on PAMAM chemistry) on the adhesion of epoxies and PU adhesives using various substrates, and as surface treatments for fibers in an epoxy matrix.

15.2 Experimental 15.2.1 Bulk Hyperbranch Incorporation

Two typical HB products were used. These products (IB-100 and AD-102) possess different MWs, degrees of functionality and viscosities (Table 15.1). Structural epoxy adhesives based on DGEBA (diglycidyl ether of bisphenol A) were cured with poly(amido-amine) and HB products at room temperature for seven days. The polyols and isocyanates used in the course of the investigation are listed in Table 15.2. The methodology used was based on partial replacement of triol (Baycoll BT-1380) with the hyperbranched component in the range 0.5–5 wt.%. The formulations were characterized using bulk and adhesion specimens. The shear adhesion strength of the PU formulations was tested by single lap shear joints, according to ASTM D 1002-94 (at a rate of 1.3 mm min–1). Peel strength was evaluated by T-Peel (ASTM 1876-95) (at a rate of 254 mm min–1).

Table 15.1 Hyperbranched polyamide products.

Property

IB-100

AD-102

Molecular weight [g mol–1] H-functionality [mol–1] Viscosity at 25 8C [Pa s] H-equivalent [g]

6500 30 130 215

12 100 45 240 270

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15 Novel Adhesion Promoters Based on Hyperbranched Polymers Table 15.2 Polyols and isocyanates used in the study.

Component

Manufacturer

Content [w/w]

Chemical composition

Baycoll BT-1380

Bayer (Germany)

50

Branched polypropylene ether triol

Baycoll BD-1110

Bayer

50

Linear polypropylene ether diol

Desmodur VK-10

Bayer

66

Mixture of 4,4'-diphenylmethane diisocyanate (MDI) isomers and higher homologues (PMDI)

EN-4/EN-7

Conap, USA

100 : 17.5

TDI-based polyurethane

The adhesive thickness was about 0.1 mm. The dimensions of the bonded area were 6 in ´ 1 in (15 cm ´ 2.5 cm). All adherends (for shear and peel) were aluminum 2024-T3 alloy pretreated by unsealed chromic anodization. The substrates were cleaned with trichloroethylene prior to bonding. The shear and peel tests were conducted on an Instron 4481 tester. 15.2.2 HB Polymers as Adhesion Promoters

Five different HB PAMAMs were included in the study: · Polyesteramide (Hyperbrane HA-1300; DSM Research, The Netherlands) [17, 18]. This is based on 15 hydroxyl end groups and has molecular weight Mn = 1300 g mol–1, Tg = –3 8C. It is soluble in ethanol. Three concentrations were used: 0.5, 1.0, and 1.5% w/w in ethanol. Each of the solutions was applied by brushing it on the adherend, dried for 30 min at room temperature and treated at 90 8C for 30 min. The adhesives included epoxy (Epon 815C; Miller Stephenson) cured with linear amido-amine (Versamide 140; Shell), and PU (EN-4/EN-7; Conap, USA). · Two HB PAMAMs were tested as primers: The first (AD-102) was chosen because primary amines react with both epoxies and PUs. The second (IB-100) has a lower functionality and lower MW (Table 15.1). For both HB polymers, seven different concentrations were used: 0.5, 1.0, 1.5, 2.0, 3.0, 4.0, and 5.0% w/w in ethanol. · Two PAMAM dendrimers were also tested: generation 3 and 4 products (G3 and G4 respectively; Aldrich). Two concentrations were used: 0.05 and 0.1% w/w in methanol. Each of the solutions was applied by brushing on the adherend, dried for 30 min at room temperature, and treated at 90 8C for 30 min. The adherends used were: · aluminum 2024 T3, bare or chromic acid anodized (no sealing) · polyetherimide (Ultem 1000; GE) · magnesium (AZ-90 alloy; Ortal Ltd.).


In some cases a silane (A-187; Union Carbide) was applied to the Al adherends as a coupling agent for better chemical bonding. Single lap shear (SLS) joints were primed with all four PAMAMs, bonded with epoxy or PU adhesives, and tested in shear according to ASTM D-1002 at a loading speed of 2 mm min–1. Scanning electron microscopy (SEM) was used to investigate the morphology of the primed adherends and the fracture surfaces of the bonded joints. The durability of joints with adherends treated with optimum concentrations of the HB polymer primer and bonded with epoxy or PU adhesives was tested on wedge joints (ASTM D-3762) with aging conditions of 50 8C/95% RH, for 1, 8, 24, and 168 h. The interfacial adhesion effect of the HB PAMAMs on fibers in epoxy composite systems was investigated with laminates prepared by compression molding. The reinforcing fabrics used were aramid (Kevlar; DuPont) and poly(p-phenylene-2,6-benzobisoxazole) (PBO) (Zylon; Toyobo). The fabrics were first treated with a 1.5% solution of HB PAMAM (AD-102), dried, then interleaved with an epoxy film structural adhesive (epoxy stage B on polyester net), and compression-molded for 90 min at 120 8C. For the aramid-based composite, FM-73 (Cytec) was used as the matrix. For the PBO-based composite, AF-191 (3M) was used as the matrix. The laminates were cut into strips and tested for interlaminar shear strength (ASTM D-2344) using a three-point bending instrument. The treated composites were compared with the same laminates but not including the surface treatment with the HB PAMAM.


The lap shear and T-peel strength results obtained for PU formulations of Baycoll BT-1380/Baycoll BD-1110/Desmodur VK-10 (50 : 50 : 66 by wt.) containing various amounts of AD-102 or IB-100 are given in Table 15.3. It can be seen that the maximum shear and peel strengths was obtained with HB polymer Epox AD-102 at a concentration of 1%. Higher concentrations resulted in decreased shear and peel strengths. The same result was obtained with HB Epox IB-100. Elevated temperatures accelerate the diffusion of the bulky HB polymer. Furthermore, the functionality of the HB polymer plays a major role in the crosslink density obtained. The lower MW polymer, Epox IB-100, has a lower functionality than the higher MW one, Epox AD-102. Structural adhesives having high shear and peel strengths in combination with a Tg of 80 8C were obtained using these novel cost-effective curing system that contained HB PAMAM. The simultaneous increase of shear and peel strength is very rare in adhesive systems. In most cases an increase in shear strength (higher crosslinking density) is accompanied by a decrease in peel strength (lower toughness). The former unique phenomenon is attributed to the incorporation of the 3D HB PAMAM that imparted 3D crosslinked network

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architecture as well as an energy-absorbing structure to the epoxy and PU adhesives. The hyperbranched HA-1300 contains hydroxyl end groups; thus the crosslinking with epoxy and PU is very limited and the wetting is inhibited. It is used as a primer. Mechanical results show that at optimal concentration (0.5 wt.%) of HA-1300 the adhesion strength is increased by only 20–33% for the epoxy adhesive (Epon 815C/Versamide 140) and not at all for PU adhesive (Table 15.3). Tables 15.4 and 15.5 present the adhesion shear strength of epoxy (Epon 815C/Versamide 140)-bonded aluminum joints (anodized) and of PU (EN-4/EN-7)-bonded aluminum joints (anodized), respectively. Both types of adhesively bonded joints were primed with high MW HB PAMAM (AD-102). As is evident from the tables, significant improvements in shear strength were ob-

Table 15.3 Shear strength [MPa] and peel strength [N mm–1] of PU BT-1380/BD-1110/VK-10 with Epox AD-102 and IB-100 formulations at various concentrations. a)

Epox AD-102 concn. [wt.%] 1 week at RT + 14 h at 80 8C 1 week at RT + 14 h at 80 8C

Peel Shear

Epox IB-100 concn. [wt.%] 1 week at RT + 14 h at 80 8C a)

Shear

0

1

2

3

4

1.7 5.4

4.4 7.9

2.6 7.8

– 7.1

– 6.5

0

1

2

3

4

4.7

7.5

7 .1

–

7.1

Standard deviation ± 5%; locus of failure was cohesive.

Table 15.4 Single lap shear strength [MPa] of epoxy-bonded Al adherends primed with various concentrations of HMW HB PAMAM (AD-102).

PAMAM Concn. [wt.%]

No treatment

No treatment with A-187

Chromic anodized (no seal)

Chromic anodized (no seal) + A-187

0 (Ref.) 0.5 1.0 1.5 2.0 3.0 4.0 5.0

4.2 ± 0.6 (A) a) 3.6 ± 1 (A/DIV) 4.9 ± 1 (A) 4.4 ± 0.7 (A) 4.6 ± 1.3 (A) – – –

7.0 ± 0.8 (A/DIV) 4.0 ± 0.9 (A) 4.9 ± 0.6 (A/DIV) 5.8 ±0.6 (C) 5.1 ± 0.7 (A/DIV) – – –

12.6 ± 0.9 (A/DIV) 12.2 ± 1 (A/DIV) 15.2 ± 1.7 (M) 16.8 ± 0.9 (M/C) 22.0 ± 2.6 (C) 17.3 ± 2.1 (C) 15.9 ± 2.6 (M/C) 14.9 ± 2.0 (M/C)

14.8 ± 1.6 (A/M) 12.8 ± 1 (A/DIV) 15.4 ± 1.1 (M) 16.9 ± 0.8 (M/C) 22.4 ± 2.4 (C) 17.6 ± 1 (C) 15.8 ± 0.9 (M/C) 13.7 ± 1 (M)

a)

Failure mode: C: cohesive; M: mixed; A: interfacial; A/DIV: interfacial divided (adhesive divided between both adherends).

15.3 Results and Discussion Table 15.5 Single lap shear strength [MPa] of polyurethane EN-4/ EN-7-bonded Al adherends primed with HMW HB PAMAM.

PAMAM Concn. [wt.%]

No treatment

0 (Ref.) 0.5 1.0 1.5 2.0 3.0 4.0 5.0

5.1 ± 0.3 5.8 ± 0.9 6.2 ± 0.4 4.4 ± 0.4 3.5 ± 0.4 – – –

a)

(A/DIV) a) (A/DIV) (A/DIV) (A) (A)

Chromic anodized (no seal) 5.8 ± 0.6 (C) 7.0 ± 1.1 (M/C) 9.8 ± 1 (C) 8.4 ± 1.2 (M/C) 9.3 ± 1.4 (M/C) 7.2 ± 0.5 (A/DIV) 8.8 ± 0.6 (C) 8.5 ± 1 (C)

Failure mode: C: cohesive; M: mixed; A: interfacial; A/DIV: interfacial divided.

tained by the HB polymer priming of the aluminum adherends. For the epoxy adhesive the optimal concentration of the primer is 2% while for PU it is 1%. SEM results show that the HMW PAMAM forms a very thin uniform coating on the adherend. High magnification reveals 1 lm uniform structures (Fig. 15.1). Micrographs of the failure surfaces show good adhesion between the adhesive and the primed surface. Fig. 15.2 shows the failure surfaces of Al adherend bonded with the epoxy adhesive. The failure mode is cohesive in the epoxy adhesive. The improved interlocking of the adhesive in the primed surface topography is presented in Fig. 15.3. The same effect as on aluminum adherend was also found with magnesium and PEI adherend. The optimum concentrations of each of the primers studied were applied to PEI (Ultem 1000) and magnesium alloy (AZ-91). Table 15.6 summarizes the single lap shear strengths of the resulting bonded joints. The results show that PAMAMs are effective in improving lap shear strength on

Fig. 15.1 Typical structures of PAMAM on aluminum adherend (X 2000). The bar is 5 lm.

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15 Novel Adhesion Promoters Based on Hyperbranched Polymers Fig. 15.2 Failure surface of 2% PAMAM-primed sample (mixed/cohesive failure) bonded with epoxy adhesive X250. The bar is 500 lm.

(bar)

Fig. 15.3 Failure surface of 2% PAMAM-primed sample (mixed/adhesive failure) bonded with PU adhesive X3300. The bar is 2 lm.

Table 15.6 Single lap shear strength [MPa] for various adherends and PAMAMs.

Primer (optimum concn.)

Adherend PEI/PEI

Mg/Mg

Al/Al

Reference (0%)

6.82 ± 1.5 (A/DIV) a)

7.44 ± 1.5 (A/DIV)

12.6 ± 0.9 (A/DIV)

HMW PAMAM (2%)

8.75 ± 0.4 (M)

10.8 ± 0.7 (C)

22.0 ± 2.6 (A)

LMW PAMAM (1.5%)

11.2 ± 0.4 (C)

8.95 ± 0.7 (C/M)

19.0 ± 2 (A)

PAMAM G4 (0.1%)

9.28 ± 0.7 (M)

7.75 ± 1.2 (M)

16.6 ± 1.2 (A)

a)

Failure mode: C: cohesive; M: mixed; A: interfacial; A/DIV: interfacial divided. Each result is an average of five samples.

15.3 Results and Discussion Table 15.7 Single lap shear strength [MPa] of epoxy-bonded Al adherends primed with various concentrations of G3 and G4.

Dendrimer

Concn. [wt.%]

No treatment

Chromic anodized (no seal)

0 (Ref.) G-3 G-3 G-4 G-4

0 0.05 0.1 0.05 0.1

7.0 ± 0.8 (A/DIV) a) 3.1 ± 0.5 (A/DIV) 3.0 ± 0.3 (A/DIV) 2.5 ±0.5 (A/DIV) 3.0 ± 0.6 (A/DIV)

12.6 ± 0.9 16.8 ± 1.5 18.0 ± 3.8 19.3 ± 2.5 19.0 ± 2.2

a)

(A/DIV) (C/M) (C) (M/C) (M/C)

Failure mode: C: cohesive; M: mixed; A: interfacial; A/DIV: interfacial divided.

both Mg and PEI. The best results for Mg were obtained in the case of the HMW HB PAMAM, which improves shear strength to 149%. On PEI, the low molecular weight (LMW) HB PAMAM improves shear strength to 164%. Table 15.7 presents the single lap shear strength of Al adherends primed with G3 and G4 dendrimers bonded with epoxy adhesive. On anodized Al, results show improvement of 33–53% in lap shear strength and in failure mode using dendrimers as a primer. The effect of the HBPs and dendrimer PAMAMs on the strength characteristics of aluminum-, PEI-, and magnesium-bonded joints based on epoxy and PU adhesives is very significant. This result can be attributed to the chemical interactions between the peripheral end groups of the branched PAMAMs and the epoxies and PUs at the joint interface. Furthermore, the interactions of the branched PAMAMs with the metallic or organic substrates may be physical or chemical in nature. Finally, the amount of the branched PAMAMs at the interface may not exceed a certain threshold, since plasticization of the epoxy and high primer thickness may cause inferior interfacial strength. For further study of the durability of the bonded joints, wedge-type specimens 6 in ´ 1 in (15 cm ´ 2.5 cm) were bonded for the formulations exhibiting the best results in shear – 2% of AD-102 for epoxy- and 1% for PU-primed joints. The wedge specimens were characterized with respect to crack growth. Crack length was measured as a function of time at 95% RH, 50 8C. Figs. 15.4 and 15.5 present the wedge test results for epoxy and PU, respectively. A rigorous examination indicates that the initial crack is interfacial and turns cohesive. In the case of the epoxy adhesive joints, the crack penetrates under the anodization layer. As can be learned from Figs. 15.4 and 15.5, the effect of the HB PAMAM primer is highly significant, leading to a very durable interface. The durability in the case of the epoxy system is improved to a higher level than for the PU system. Comparison of the results for the HMW PAMAM with those for the LMW PAMAM leads to the conclusion that the optimal concentration of the HB PAMAM that yields the highest shear strength depends not only on the adhesive used but also on the MW and the functionality of the primer. For the epoxy ad-

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15 Novel Adhesion Promoters Based on Hyperbranched Polymers Fig. 15.4 Crack propagation in epoxy-bonded wedge specimen treated with 2 wt.% AD-102.

hesive, the optimal concentration of the LMW HB PAMAM primer is 1.5 wt.% while for the HMW HB PAMAM it is 2 wt.%. The mechanical results of interlaminar shear strength for the two kinds of laminates prepared using hyperbranches impregnated in the reinforced fibers are presented in Table 15.8. The two laminates tested included Kevlar fabric primed with 1.5 wt.% HMW PAMAM (AD-102) and bonded with FM-73, and Zylon fabric primed with 1.5 wt.% HMW PAMAM (AD-102) and bonded with AF-191. The results show that composites made of primed HMW HB PAMAM aramid and PBI fibers in epoxy matrices demonstrate much improved interlaminar shear strength (the Kevlar laminate shear strength increased to 160% and that of the Zylon to 140%).

Fig. 15.5 Crack propagation in PUbonded wedge specimen treated with 1 wt.% AD-102.

15.4 Conclusion Table 15.8 Interlaminar shear strength (ASTM D-2344) for laminates with and without HMW PAMAM.

Laminate

Interlaminar shear strength [MPa]

Aramid (Kevlar) Untreated AD 102

11.4 ± 0.5 18.1 ± 1.9

PBO (Zylon) Untreated AD 102

9.0 ± 0.8 12.6 ± 1.7

15.4 Conclusion

The use of hyperbranches as primers or as crosslinking agents for the epoxy and PU adhesives resulted in improved shear strength and higher adhesion durability (wedge test results). The modification of PU and Epoxy systems with the hyperbranched or dendritic polymers was effective using small amounts. HB polymers were more effective than the dendrimers for shear strength improvement. Epoxy adhesive showed better improvement with all HB polymer and dendrimer primers. The optimum concentration of primer solution depends on the kind of adhesive and on the kind of primer. This proves that the main mechanism of improvement is a chemical interaction. For epoxy adhesive the optimum amount of AD-102 was 2 wt.% and for IB100 it was 1.5 wt.%. For PU, the optimum amount of AD-102 was 1% (the shear strength improved for both treated and untreated Al substrates). Dendrimers had a lower improvement effect than hyperbranches but the concentration needed for the same effect was an order of magnitude lower. Hyperbranches showed shear strength promotion on various adherends – plastic and metal. Morphological observation by SEM showed interlocking of the adhesive in the adherend surface due to microstructures (star-like or spheres) formed by the hyperbranches or dendrimer clusters on the interface. Very significant improvement in interlaminar shear strength was observed for the Kevlar/epoxy and Zylon/epoxy composites with fibers treated with hyperbranch AD-102. The 3D architecture obtained by the introduction of 3D molecules should lead to new approaches to enhancing simultaneously the shear and peel properties of polyurethane adhesives and coatings, which otherwise is difficult to achieve.

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References 1 D. A. Tomalia, Polymer 1985, 17(1), 117. 2 A. Hult, M. Johansson, E. Malmstrom, 3 4

5

6

7 8

Adv. Polym. Sci. 1998, 143, 2–34. E. Malastrom, A. Hult, Rev. Macromol.Chem. Phys. 1997, C37, 555–579. L. Moshinsky, S. Kenig, H. Dodiuk-Kenig, A. Buchman, “Epoxy Adhesives and Primers Based on Hyperbranched Polyamidoamines”, Meeting of the Adhesion Society, Williamsburg, VA, Feb. 2000. O. A. Matthews, A. N. Shipway, J. J. Stoddart, “Dendrimers – Branching out from Curiosities into New Technologies”, Prog. Polym. Sci. 1998, 23(1), 1. G. R. Newcome, C. N. Moorefield, F. Voegtle, “Dendritic Molecules, Concepts, Syntheses and Perspective”, VCH, Weinheim, 1996, p. 261. Eur. Patent Appl. EP 66 366 (1982). D. A. Tomalia, V. Berry, M. Hall, D. M. Hedstrand, Macromolecules 1987, 20(5), 1164.

9 D. A. Tomalia, W. A. Goddard, Angew.

Chem. 1990, 102(2), 119. 10 US Patent US 5 760 142 (1998). 11 A. J. Allen, Hercules Inc., Eur. Patent.

Appl. EP 802 215 (1997). 12 Tomalia, D. A., Dow Chemical Co., US

Patent US 4 737 550 (1988). 13 Tomalia, D. A., Dow Chemical Co., US

Patent US 4 631 337 (1986). 14 Tomalia, D. A., Dow Chemical Co., US

Patent US 4 558 120 (1985). 15 Israel Patent Appl. 125 565 (1998). 16 US Patent Appl. 09/295 320 (1999). 17 PCT/IL Patent 99/00540, “Highly

Branched Oligomers, Process for their Preparation and Applications Thereof” (1999). TM 18 Hyperbane DSM New Business Development, 99-36 (1999).

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16 Rheology of Hot-Melt PSAs: Influence of Polymer Structure C. Derail and G. Marin

Abstract

At room temperature, the hot-melt pressure-sensitive adhesives (HMPSAs) satisfy the Dahlquist criterion, and consequently have permanent tack. It is well established that the tack properties of these adhesives are governed, to a large extent, by their rheological properties. We have studied the full formulations based on triblock and diblock copolymers, and also on new molecules, such as tetrablock or radial copolymers, designed to improve the end-user properties. We describe the dynamic mechanical properties in the linear domain, at room temperature, of the pure copolymer blends and the full HMPSA formulations. We propose a model, based on molecular dynamics concepts, which describes the rheological behavior in a very wide range of frequencies, for all the copolymers and full formulations in the study. The effect of the diblock copolymer content on the rheological behavior of the adhesive and the link with the morphology of the copolymers is discussed in detail. Finally, we conclude that, in using molecular dynamics concepts, it is possible to propose an efficient “virtual formulation” tool for the diblock–triblock systems studied.

16.1 Introduction

The relationship between the rheological properties and the adherence properties of soft polymer-based adhesives has been widely studied, particularly since the pioneering work of Gent and Petrich [1]. The design of adhesive formulations is still mostly based on the intuition and experience of the formulator for most industrial applications using hot-melt adhesives. Recent research results show, however, that it is possible to establish strong quantitative links between the molecular parameters of the components of a given formulation and its adherence properties through its rheological behavior [2–7]. Based on the work of Zosel (who was the first to use optical observations during the debonding proAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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16 Rheology of Hot-Melt PSAs: Influence of Polymer Structure

cess [8]) on the measurement of adherence properties using the probe–tack test [9], it has been shown recently that one may establish a clear qualitative link between the nature of the copolymer-based adhesive and the nature of the debonding process [10, 11]. Hot-melt pressure-sensitive adhesives (HMPSAs) are generally used at room temperature, where they have to exhibit aggressive tack properties. HMPSAs classically used for label and tape applications are basically made of a copolymer base and tackifying resins. Usually, the coating process is performed in the melt state at high temperature. In such systems, it is well established that the copolymer brings cohesion to the adhesive and the resin generally shifts the glass transition temperature and gives tack properties to the formulation [5]. It is important that the full formulation must exhibit a solid-like viscoelastic behavior in order to obtain the expected tack properties (see the Dahlquist criterion [3]) and no creep behavior at long times. These features may be obtained through ordered morphologies, as in the case of block copolymers. We will focus in this paper on the rheological properties, at room temperature, of styrene–isoprene block copolymers, particularly Triblock [SIS]–Diblock [SI] copolymer blends. We will describe the effect of the molecular parameters of the copolymers on the rheological behavior, and will propose, on the basis of molecular dynamics models derived from the reptation concept and the analysis of the dynamic behavior of the blend [SIS–SI], a model which allows calculation of the variation of the complex shear modulus as a function of frequency. Different types of macromolecules have been designed from calculations using this molecular model in order to improve the processing and end-user properties of the full formulations (HMPSAs). In Section 16.2 we will describe the main features expected from a “good adhesive” and we will present the rheological behavior of pure copolymers ([SIS] and [SI]) and blends ([SIS–SI]) as well as full adhesive formulations ([SIS–SI–resin]). In Section 16.3, we will focus on the rheological model which is based on concepts of molecular dynamics and describes reasonably the viscoelastic properties of the pure blends and the full formulations. In Section 16.4, we will describe how we have used a predictive approach to “calculate” the characteristics of new molecules. We compare the rheological properties of “regular” [SIS–SI] copolymer blends with the analogous behavior of these newly designed molecules at room temperature. The comparison is also presented for the full formulations. We will focus on the variations of the secondary rubbery plateau as a function of polyisoprene content. Finally (Section 16.5), we will conclude on the improvement of formulations by using the model which then becomes a tool for “virtual” formulation.

16.2 Main Features of the Viscoelastic Behavior of the Pure Components

16.2 Main Features of the Viscoelastic Behavior of the Pure Components, Blends, and Full Adhesive Formulations

The HMPSAs (full adhesive formulations) [SIS–SI–resin] presented in this work are formulated with: · a pure blend based on [SI] diblock and [SIS] triblock copolymers. The diblock content can be varied from 30% to 90%. In this study, the proportion of total copolymer blend is fixed (31% of the full formulation). · tackifying resins (Escorez 1310, manufactured by ExxonMobil Chemical Co. (Houston, TX, USA) and Wingtack 10, manufactured by the Chemical Division of the Goodyear Tire and Rubber Company (Akron, OH, USA)). The proportion of total resin in the full formulation is 69%. 16.2.1 Rheological Experiments

The pure copolymers, pure resins, pure blends [SIS–SI], and formulations have been characterized rheologically by measuring the complex shear modulus (G' and G'') as a function of circular frequency, x, at various temperatures (from – 60 8C to +30 8C). These mechanical spectroscopy experiments were performed in the frequency range 10–2–100 rad s–1 using a Rheometric RDA II rotational rheometer in parallel-plate geometry. The time–temperature superposition principle can be used to plot the main relaxation domains from the terminal zone region (flow region or secondary plateau region for viscoelastic solids like [SIS]) up to the glassy domain (Ta relaxation) at high frequencies. In this article, the master curves have been reported at Tref = 20 8C. We have verified that the morphology of the copolymers is not affected by the variation of the temperature and it is important to note that the maximum temperature for the experiments performed on copolymers was 30 8C. 16.2.2 Rheological Behavior of the Pure Components: [SI], [SIS], and Pure Blends

Fig. 16.1 provides the master curves for the rheological behavior of the pure components [SI] and [SIS] (Mw = 107 000 g mol–1, 16.5% S), [SIS] (Mw = 118 000 g mol–1, 18% S) at room temperature compared with the rheological behavior of a pure entangled polyisoprene (Mw = 29 000 g mol–1). The styrene content is reported in weight %. One can observe various regions corresponding to the usual relaxation domains, ranging from the glassy zone to the terminal region, which are typical of the viscoelastic behavior of entangled polymer melts. For the [SI] diblock and the [SIS] triblock copolymers, one can observe that the glassy domain (with the same value of frequency at the maximum of G'', for example) is about the same as for the pure polyisoprene (implying the same apparent a transition) whereas, in the terminal domain, one can notice a

231

Fig. 16.1 Master curves, Tref = 20 8C, of storage and loss moduli versus frequency for [SI] (Mw = 107 000 g mol–1, S = 16.5%), [SIS] (Mw = 118 000 g mol–1, S = 18%) and pure polyisoprene (Mw = 29 000 g mol–1). Full lines: model. The different domains are indicated for each polymer and copolymer. `, G', n, G'' for [SI]; ~, G'; s, G'' for [SIS]; *, G', *, G'' for I.

232


16.2 Main Features of the Viscoelastic Behavior of the Pure Components

very different type of behavior. The triblock copolymer exhibits an elastic plateau region without a cross-over between G' and G'', indicating a solid-like behavior which is independent of the molecular weight of the triblock copolymer. For the diblock copolymer, one can observe a flow region at intermediate frequencies and, at the lowest frequencies, the beginning of a secondary elastic plateau due to the organization of the polystyrene sequences in the polyisoprene matrix. The copolymer base of the adhesives is a blend of diblock and triblock copolymers. In Fig. 16.2 we compare the rheological behavior for two blends which have different diblock contents (54% and 85%), different molecular weights for the diblock part (86 000 and 104 300 g mol–1), and the same styrene content (16%). The molecular weight of the triblock is 128 000 g mol–1. Let us describe now the relaxation domains exhibited by the two blends: · a transition region from the glassy domain down to the rubbery domain (a relaxation); · a rubbery plateau domain; · a decrease in the storage modulus corresponding to the relaxation of the isoprene sequence of the diblock copolymers: when the molecular weight of the isoprene part of the diblock increases, the characteristic frequency of this relaxation domain decreases (see the shift of the maximum of tan d); · a terminal zone corresponding to the secondary plateau modulus (Gs) of the triblock part swollen by the isoprene sequence of the diblock: in this zone, the level of the plateau decreases when the diblock content increases. This last domain gives a specific viscoelastic solid behavior to the blend. As a consequence, we have particularly studied the behavior of this terminal region. The glassy domain (higher frequencies) is not reported in Fig. 16.2 because we focus our discussion on the terminal region. 16.2.3 Rheological Behavior of the Full Adhesive Formulations

The pure diblock–triblock blends previously described cannot develop sufficient tack onto surfaces, in particular because the modulus level in the terminal zone is too high (Dahlquist criterion [3]). In order to improve tack properties, tackifying resins are added to the copolymer blend. They have in fact two important effects [4, 5]: · a thermodynamic effect which generally shifts the glass transition temperature to temperatures closer to the end-use temperature · a topological effect which decreases the elastic modulus, in particular for the secondary elastic plateau. These two effects bring aggressive tack properties to the copolymer base. One can see an example for a classical full adhesive formulation in Fig. 16.3. The different relaxation domains previously described are the same. It is important to notice that, in the terminal region, as for the pure blend, a low secondary pla-

233

Fig. 16.2 Master curves, Tref = 20 8C, of storage and loss moduli and tan d versus frequency for [SIS–SI]1 (SI = 85%, Mw of the SI = 104 300 g mol–1, S = 16%) and [SIS–SI]2 (SI = 54%, Mw of the SI = 86 000 g mol–1, S =16%). Mw of the SIS = 128 000 g mol–1 in both cases. Full lines: model. The different domains are indicated for [SIS–SI]1. `, G', n, G''; - - - -, tan d for [SIS–SI]1; ~, G'; s, G''; ——, tan d for [SIS–SI]2.

234


Fig. 16.3 Master curves, Tref = 20 8C, of storage and loss moduli versus frequency for [SIS–SI–resin] (SI = 85%, Mw of the SI = 104 300 g mol–1, Mw of the SIS = 128 000 g mol–1). Full lines: model. `, G', n, G''.

16.2 Main Features of the Viscoelastic Behavior of the Pure Components 235

236


teau modulus brings a no-flow behavior as well as relevant elastic properties for a good tack. In Section 16.4, we will particularly discuss the variations of the secondary plateau as a function of the diblock content in the blend [SIS–SI] and in the full adhesive formulation. The secondary plateau value is really a key parameter for optimizing the rheological behavior of the final adhesives.

16.3 A Model of the Rheological Behavior 16.3.1 A Model for the Pure Copolymers

The linear viscoelastic behavior of the pure polymer and blends has already been described quantitatively by using models of molecular dynamics based on the reptation concept [12]. To describe the rheological behavior of the copolymers in this study, we have selected and extended the analytical approach of Benallal et al. [13], who describe the relaxation function G(t) of linear homopolymer melts as the sum of four independent relaxation processes [Eq. (1)]. Each term describes the relaxation domains extending from the lowest frequencies (Gc(t)) to the highest frequencies (GHF(t)), and is well defined for homopolymers in Ref. [13]. Gt Gc t GA t GB t GHF t

1

We recall below the relevant expressions for homopolymers and how we have extended this approach to copolymers. In all the expressions given below, s0 is the elementary relaxation time of the Kuhn segment, N is the number of entanglements of a given copolymer sequence or macromolecular chain, Ne is the number of monomers between entanglement points (Me being the molecular weight between entanglements), and Na is the number of monomers per chain (or sequence) for the pure polyisoprene. Let us describe the different relaxation domains for the copolymers in the study: · the a relaxation (glass transition) at very short times where only local relaxation motions are observed. We have used a Davidson–Cole equation with a dispersion exponent of 0.5 [14]. The relevant relaxation modulus, GHF (HF = high frequencies) is given by Eq. (2), where the two parameters G! and s0HF used for the copolymers are the same as for the a relaxation of a pure polyisoprene homopolymer.

16.3 A Model of the Rheological Behavior

GHF t G1 exp

t

0:5

s0HF

2

· two Rouse relaxation processes at intermediate times. In this intermediate domain of frequencies, the relevant relaxation moduli can be expressed by Eqs. (3, 4) [15], where the characteristic times for the A and B processes are given by Eqs. (5) and (6) respectively: G0n can be calculated by using the rubber elasticity equation [16]. For the copolymers these expressions will be the same as for the homopolymer. GA t G0n

GB t G0n

2 pt exp s A p1

3

2 Ne X 1 pt exp N s B p1

4

Ne X

1 sA s0 Ne2 6

5

1 sB s0 Na2 3

6

· the terminal relaxation process at long times (process C), which is the signature of reptation for the homopolymers. We recall in this case the relevant equation in the terminal domain, Eq. (7), where sc, the reptation time, is given by Eq. (8). GC t

G0n

2 X 8 pt exp 2 p2 p s c p

sC s0 N 3 Ne2

7 8

We have seen previously that the diblock and triblock copolymers do not exhibit reptation in the terminal domain. Therefore, the terminal relaxation process in Benallal’s model has been modified in the case of diblock and triblock copolymers and their blends: · For the diblock copolymer, which exhibits a flow region at longer times than pure polyisoprene, the relaxation of the isoprene sequence is treated like the relaxation of the arm of a star polymer. We have followed the description proposed by McLeish [17, 18] for star homopolymers. The distribution of relaxation times is given by Eq. (9), where Mb is the molecular weight of one branch (here the molecular weight of the polyisoprene sequence), s ranges be-

237

238


tween 0 and Mb and m', a numerical parameter linked to the structure of the network [17], has a value close to 0.6. sstar s

s00

2 6m0 s exp Me Mb 2

s3 3Mb

9

The elementary time s00 is a function of s0 [Eq. (10)]. s00

4 M2 s0 b2 3 M0

s pMe m0 Mb

10

In this domain, the expression of the relaxation modulus is Eq. (11), which also leads to a distribution of relaxation times which is much broader than the spectrum of linear species. 2G0n Gstar t Mb

ZMb 1 0

s t exp ds Mb ts

11

· A phenomenological approach has been proposed to describe the effect of the organization of the copolymers on the rheological behavior in the low-frequency range [19]. Following this approach, we propose, for the [SI] diblock and [SIS] triblock copolymers, a power law according to Eq. (12), where K is a constant and n is an exponent which can be varied with respect to the organization of the copolymer. In our case, the value of n is kept constant for all copolymers and full formulations (n = 0.1). Gs t K tn

12

This “critical gel” behavior describes reasonably the terminal relaxation domain of pure diblock and triblock copolymers. Lastly, the overall relaxation moduli of the diblock and the triblock copolymers are expressed in Eqs. (13, 14) as the sum of the individual contributions given above. The factor f(Usty) takes into account the filler effect of polystyrene [20– 22]. Gdiblock t GHF t f Usty GA t GB t Gstar t Gs t

13

Gtriblock t GHF t f Usty GA t Gs t

14

A Fourier transform makes it possible to shift from the time domain to the frequency domain and yields the complex shear modulus for the diblock. In Fig. 16.1, one can observe the good agreement between experimental data and calculations for the pure diblock and triblock copolymers.

16.3 A Model of the Rheological Behavior

16.3.2 A Model for the Blends [SIS–SI]

We have proposed the use of a quadratic blending law of the “double reptation” type to express the viscoelastic behavior of [SIS–SI] blends based on the viscoelastic behavior of the diblock and the Utriblock copolymers. It may be expressed as Eq. (15), where triblock is the volume fraction of the triblock copolymer in the [SIS–SI] blend. G0:5 SIS

SI t

Utriblock G0:5 triblock t 1

Utriblock G0:5 diblock t

15

We have already shown [23] that the double reptation law can also reasonably be applied to the complex shear modulus, which simplifies the calculations by avoiding Fourier transforms from the time domain to the frequency domain. In Fig. 16.2, one can observe the good agreement between experimental data and calculations for pure [SIS–SI] blends. 16.3.3 A Model for the Full Adhesive Formulations [SIS–SI–Resin]

For the full formulations, the resin is considered as an antiplastifying solvent of the isoprene matrix, which is treated as a concentrated polymer solution [24, 25]. As previously described, the resin shifts the glass transition temperature of the isoprene part. It is possible to explain/predict the shift from the master curves of the pure [SIS–SI] by calculating the modified elementary times s0 and s0HF through Eqs. (16, 17), taking into account the change in the mobility factor (friction coefficient) [5, 25] and considering an increase in the Vogel temperature (T?) of the WLF (Williams, Landel and Ferry [16]) equation. s0HF A exp 0

s0 A exp

B T0

B0 T0

16

T1

17

T1

T0 is the reference temperature (here 20 8C), A, B, A', and B' are obtained by a leastsquares method [5, 25] and T? is given by Eq. (18), where Uresin1, Uresin2, U[SIS–SI], T?,resin1, T?,resin2, T?,[SIS–SI] are the volume fractions and the Vogel temperatures of the resin parts and the copolymer part in the full adhesive formulation. T1 adhesive Uresin1 T1;resin1 Uresin2 T1;resin2 USIS

SI

T1;SIS

SI

18 It is important to point out that we have demonstrated that the two copolymers in the study, [SI] and [SIS], have, as a first approximation, the same characteris-

239

240


tic temperatures (see Fig. 16.1) as the pure polyisoprene. So the variation of the elementary times is the same in both cases. We noticed previously that, when resins are added to the copolymer blend, the plateau modulus decreases (which is what we call a “dilution effect”, as opposed to the “thermodynamic effect” on the Tg described above). The entanglement network of the isoprene part of the copolymers is swollen and the molecular weight between two entanglement points increases by following the relevant power law [Eq. (19)], where the experimental values of n may range from 2 (theoretical value) to 2.5 for concentrated polymer solutions and melts as well as for the copolymers in the present study [4, 5, 20, 24, 25]. This topological effect is the same as for the isoprene part of diblock and triblock copolymers. G0n;formulation G0n;SIS

SI

UnSIS

SI

19

One can verify (Fig. 16.3) the good agreement between experimental data and calculations for [SIS–SI–resin] full adhesive formulations. 16.4 Discussion 16.4.1 Molecular Design

In order to improve the processing and end-user properties and also to simplify the preparation of the blends described above, we propose to design new copolymers which would have the same block sequences (polyisoprene and polystyrene) but not the same topology [23, 26–28]. We have already pointed out that, to obtain optimized rheological properties for a good tack, we must have a free polyisoprene sequence which explores the network of the trapped polyisoprene sequences of the triblock copolymer and swells this network. It is important to notice that this is the triblock copolymer which causes the solid-like behavior [23]. This configuration can be obtained with a series of block sequences terminated by a polyisoprene sequence [26–28]: polystyrene–polyisoprene–polystyrene–polystyrene–(. . .)– polyisoprene. We focus in this article on the simplest configuration made of four blocks (SISI). In this case, the free polyisoprene end sequence plays the same role as the free polyisoprene part of the diblock in the [SIS–SI] blends, and the isoprene part trapped between the two sequences of polystyrene plays the same role as the isoprene sequence of the triblock copolymer. The characteristics of this tetrablock must be submited to the same constraints as one which are described for the [SIS–SI] blends: · The styrene content must be the same as the average content in [SIS–SI] blends. · The molecular weight of the end polyisoprene (free sequence) must be the same as for the free polyisoprene sequence of the SI diblock.

16.4 Discussion

· The molecular weight of the trapped polyisoprene sequence must be higher than the critical molecular weight for entanglements of polyisoprene. In the same way, a blend with a star copolymer with n branches of SI (the polystyrene sequences being outside the star), blended with free diblock copolymer and respecting the constraints developed for the tetrablock, is a good candidate to replace the [SIS–SI] blends. We have calculated from the model the characteristics of these newly designed structures, which have been synthesized [23], and we have measured their rheological properties at room temperature. As Fig. 16.4 shows, it is possible to mimic the rheological behavior of the [SIS–SI] blend (we have applied a vertical shift on the curves for greater clarity). Also, the rheological behavior of the full formulations (Fig. 16.4) based on these newly designed copolymers yields the same properties as for [SIS–SI], which demonstrates that the tackifying resins act like a solvent of the elastomeric part of all the copolymers presented here. Furthermore, the modeling described previously can be applied to the newly designed copolymers, as shown in Fig. 16.4 for the star copolymer. The design of new architectures represents a new way to improve the enduser properties of HMPSAs based on block copolymers.

16.4.2 On the Variation of the Secondary Elastic Plateau Modulus

We have particularly studied the variation of the secondary elastic plateau modulus (Gs) of the pure [SIS–SI] and formulated [SIS–SI–resin] blends because this is a key parameter for an optimized formulation. In this work Gs is the value of G' at a fixed frequency (10–3 rad s–1). Fig. 16.5 shows the variation of Gs for all the pure copolymers in the study [20, 23] – [SIS–SI], [SISI] and [SIn–SI] – as a function of the polyisoprene content in the triblock part or equivalent triblock part (for the tetrablock and radial copolymers). Equation (20), where UI=SIS is the volume fraction of the polyisoprene in the SIS copolymer of the [SIS–SI] blends, describes this variation at diblock contents lower than 70%. 0 0 x ! 0 GSIS x ! 0 U2I=SIS GS GSIS=SI

20

The straight line in Fig. 16.5 shows that Eq. (20) can be applied at low diblock contents. At high diblock contents (above 70%) Gs seems to tend to a limit. There is a change in behavior at 70%. This feature is more clearly seen in Fig. 16.6 for the full adhesive formulations, where we can observe two tendencies. The Gs values of the formulations based on radial copolymers tend to a limit which corresponds to the value of Gs of a pure diblock formulated with the same resins as for the blends. For the tetrablock material, the power law [Eq. (20)] still applies at high equivalent diblock contents.

241

Fig. 16.4 Master curves, Tref = 20 8C, of storage and loss moduli versus frequency for [SIS–SI] (MwSIS = 128 000 g mol–1, MwSI = 75 000 g mol–1, S = 15%), [SISI] (´ 100) (Mw = 135 000 g mol–1, S = 16%) and [SIn–SI] (´ 10 000) (Mw = 415 000 g mol–1, S = 17%) and for full adhesive formulation (right-hand scale) based on [SIn–SI] (´ 10 000). Full lines: model. `, G', n, G'' for [SIS– SI]; *, G', *, G'' for [SISI]; ~, G', s, G'' for [SIn–SI] and ^, G', ^, G'' for [SIn–SI–resin].

242


Fig. 16.5 Secondary plateau modulus for samples from Refs. [20] and [23] versus polyisoprene volume fraction in the SIS part (or equivalent) in the different copolymers. n, [SIS–SI]; *, [SIn–SI]; ^, [SISI]; *, copolymer blend with polyisoprene added.

16.4 Discussion 243

Fig. 16.6 Secondary plateau modulus for samples from Refs. [20] and [23] versus polyisoprene volume fraction in the SIS part (or equivalent) in the full adhesive formulation. *, [SIn–SI]; ^, [SISI]; *, pure formulated diblock; s, pure polyisoprene.

244


16.5 Conclusions

One can conclude that the diblock copolymer seems to “impose” its behavior on blends at high diblock content. The different behavior at high diblock contents is indeed linked to the morphology of the diblock copolymer [29]. The morphology of the block copolymers can be determined from their phase diagram. It depends on the molecular weight and ratio of the different sequences within the copolymers and on the product between the Flory interaction parameter (v) and the degree of polymerization (N) [30–32]. It is possible to measure the order–order (TO–O) or order–disorder (TO–D) temperatures by using thermomechanical analysis [29, 31–33]. The thermomechanical analysis data (2 8C min–1 and 1 rad s–1) reported in Fig. 16.7 shows that the two temperatures (TO–O = 140 8C and TO–D = 175 8C) are similar for the pure diblock and the [SIS–SI] blend with SI = 85%. This point shows that the behavior of the diblock copolymer is really a key parameter for the [SIS–SI] blends with high diblock content, by controlling their behavior at low frequencies. As far as the tetrablock copolymers are concerned, it seems on the contrary that there is no lower limit to the level of the secondary plateau. As Fig. 16.5 indeed demonstrates, the level of the secondary plateau for tetrablock materials follows the same power law variation even at high diblock contents. On the basis of these conclusions, some solutions may be proposed to correct the wrong effect of the morphology of the diblock copolymer. We therefore added a pure polyisoprene homopolymer to the blends which departed from the power law of Eq. (20). It can be observed in Fig. 16.5 that the level of the secondary elastic plateau modulus of these new blends [SIn–SI–I] [33, 34] can be corrected, and the power law [Eq. (20)] then obeyed, even at high diblock contents.

16.5 Conclusions

The understanding of the links between the microstructure and adhesive properties of copolymer blends and formulations used in HMPSA applications is important to provide quantitative tools for predictive formulation. More generally, the link between the rheological behavior of block copolymers and their phase diagram is not yet well established for blends of copolymers. In this respect, we have shown that a model based on concepts of molecular dynamics describes the overall rheological features of [SIS–SI] blends used for HMPSA applications. The model parameters are the molecular parameters of the copolymers: [SI] and [SIS] molecular weights, styrene content, and SIS/SI ratio. The model may help to design new molecules leading to an expected rheological behavior corresponding to a selected application or to given process properties. We have demonstrated that the [SI] diblock content is very important for the application, as there is a lower limit for the elasticity at high diblock content. Finally, we think that it will be possible to “play” also on the organization of the

245

Fig. 16.7 Storage modulus versus temperature for a pure diblock (Mw = 75 000 g mol–1, S = 15%) and [SIS–SI] (MwSIS = 128 000 g mol–1, MwSI = 75 000 g mol–1, S = 15%, S = 85%). `, [SI]; s, [SIS–SI].

246


References

diblock by changing the volume fraction of the two components and/or by adding pure polyisoprene or polystyrene to the blend. Using molecular dynamics concepts, we have proposed “model solutions” which lead to a reasonable description of the viscoelastic properties and especially give an efficient “virtual formulation” tool for these systems.

Acknowledgments

Part of this work has been supported by ExxonMobil Chemical Co. in the framework of a PhD program.

References 1 A. N. Gent, R. P. Petrich, Proc. R. Soc.

2 3

4 5 6 7 8 9 10 11 12 13

14

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London A: Math. Phys. Sci. 1969, 310, 433–448. A. N. Gent, J. Schultz, J. Adhesion 1972, 3, 281–294. R. L. Patrick, Treatise on Adhesion and Adhesives: Materials, Vol. 2, Marcel Dekker, New York 1969, p. 129. Ph. Vandermaesen, G. Marin, J. Adhesion 1993, 43, 1–15. C. Derail, A. Allal, G. Marin, Ph. Tordjeman, J. Adhesion 1997, 61, 123–157. M. D. Gower, R. A. Shanks, Macromol. Chem. Phys. 2004, 205, 2139–2150. A. E. O’Connor, N. Willenbacher, Int. J. Adhesion Adhesives 2004, 24, 335–346. A. Zosel, Colloid Polym. 1985, 263, 541– 553. H. Lakrout, P. Sergot, C. Creton, J. Adhesion 1999, 69, 307–359. S. Poivet, F. Nallet, C. Gay, P. Fabre, Europhys. Lett. 2003, 62(2), 244–250. A. Roos, C. Creton, Macromolecular Symposia 2004, 214(1), 147–156. P. G. de Gennes, J. Chem. Soc. 1971, 55, 572–575. A. Benallal, G. Marin, J. P. Montfort, C. Derail, Macromolecules 1993, 26, 7229– 7235. K. L. Ngai, D. J. Plazek, J. Polym. Sci., Part B: Polym. Phys. Ed. 1986, 24, 619– 632. L. Viovy, J. Polym. Sci., Part B: Polym. Phys. Ed. 1985, 23, 2423–2433.

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17 Preparation and Characterization of UV-Crosslinkable Pressure-Sensitive Adhesives H.-S. Do, S.-E. Kim, and H.-J. Kim

Abstract

UV-crosslinkable acrylic pressure-sensitive adhesives (PSAs) were synthesized by solution polymerization with 2-ethylhexyl acrylate (2-EHA), vinyl acetate (VA), acrylic acid (AA), 2-hydroxyethyl methacrylate (2-HEMA), and an unsaturated benzophenone derivative (P-36, photoinitiator), along with varying contents of 2HEMA and photoinitiator. The UV-crosslinking behavior of the PSAs was studied by FTIR-ATR spectroscopy and PSA performance was characterized by probe tack, peel strength, and shear adhesion failure temperature (SAFT) methods. Because 2-HEMA acted as a good hydrogen donor to benzophenone, its addition caused the efficiency of the photoreaction to increase, so with increasing contents of 2-HEMA and photoinitiator in the PSAs, the conversion of benzophenone groups in the PSAs at low UV doses became faster. In addition, because the crosslinking reaction proceeded via photoreaction mainly between the 2-HEMA and photoinitiator, the probe tack and peel strength of the PSAs having high concentrations of 2-HEMA and photoinitiator rapidly decreased in the early stage of UV irradiation due to increased crosslinking density. These phenomena were also observed in the SAFT test, with PSA containing high levels of 2-HEMA and photoinitiator showing high SAFT values at low UV doses.

17.1 Introduction

Pressure-sensitive adhesives (PSAs) are used in various forms for packaging, medical and masking tapes and labels. PSAs are viscoelastic materials that exhibit solidlike and liquid-like behavior so that their performance can be evaluated by tack (the capability of a PSA to adhere instantly under light pressure), peel strength (the tensile force required to remove a PSA) and shear strength (flow resistance) [1, 2]. The techniques used in manufacturing PSAs are solution and emulsion polymerization and hot-melt process. In general, acrylic PSAs are produced by soluAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

250

17 Preparation and Characterization of UV-Crosslinkable Pressure-Sensitive Adhesives

tion and emulsion polymerization due to their economic efficiency and processability, while hot-melt PSAs are made by using ABA-type styrene block copolymers [3]. PSAs made by these processes are crosslinked only by physical crosslinking, van der Waals forces, or hydrogen bonding, and they therefore exhibit limited mechanical and thermal properties. Hence, a crosslinking agent is needed to induce chemical crosslinking. When a crosslinking agent is added to PSAs, as the crosslinking density increases they become relatively nontacky materials with decreased tack and peel. Therefore, the degree of crosslinking should be carefully controlled [4, 5]. However, these systems suffer the disadvantage of being composed of two components, PSA and the crosslinking agent, and this reduces their pot life [4]. Among the crosslinking methods recently used in PSAs, radiation curing is adapted due to its economical (fast curing) and environmental (low VOC) advantages [5–8]. UV-curing techniques used in the PSA crosslinking process can be divided into UV-polymerization and UV-crosslinking methods. UV-polymerizable PSAs are formulated from oligomers, monomers, tackifiers, and photoinitiators, blended in a reactor above the Tg of the tackifier, and are then polymerized by UV irradiation after being coated on a carrier for various applications. UV-crosslinkable PSA systems are similar to UV-polymerizable PSAs for which the curing proceeds by UV irradiation, but the molecular weight of the UV-crosslinkable PSAs is higher than that of the oligomers or monomers used in UV-polymerizable PSAs and as they have UV-crosslinkable sites polymerized in the PSA backbone, they do not need further addition of photoinitiators. The curing reaction of UV-polymerizable PSAs and UV-crosslinkable PSAs is illustrated in Fig. 17.1. The main areas of research into UV-polymerizable PSAs are type, molecular weight, polarity, glass transition temperature (Tg), functionalities, and blending ratio of oligomers and monomers. Photoinitiator type (type I or type II) and kinds of tackifiers are also considered in the manufacture of PSA. As a type I

Fig. 17.1 Schematic illustration of curing of (a) UV-polymerizable PSA; (b) UV-crosslinkable PSA.

17.1 Introduction

Scheme 17.1 Radical formation of type I and type II photoinitiators.

photoinitiator is an intramolecular photocleavage type and type II photoinitiator is an intermolecular hydrogen abstraction type, when they are exposed in UV irradiation type I photoinitiator gives a benzyl radical and an alkyl radical, and type II photoinitiator, by abstracting hydrogen from a neighboring group, forms a ketyl radical (Scheme 17.1). Areas of research into UV-crosslinkable PSAs are aimed to improve the efficiency of UV-curing and the adhesion properties by controlling the molecular weight and the PSA polymer composition, and by adding additional resins. Both of these systems have the advantage of fast curing at room temperature; unreacted photoinitiator and monomer in UV-polymerizable PSAs may cause skin irritation to final users, have an unpleasant odour, and show unstable adhesion properties, but as the monomers and photoinitiator are all copolymerized in UV-crosslinkable PSAs, they do not release any residual monomer after UV irradiation. UV-curable PSAs show variations in adhesion properties as the above conditions (the kinds of monomer, oligomer, photoinitiator, and tackifier) are varied and also their properties can be changed by varying the curing conditions (UV dose, film thickness, and curing temperature). By using commercially available, UV-crosslinkable, acrylic hot-melt adhesives (acResin® A 258 UV, supplied by BASF), the peel strength test was performed with varying coating thickness [9]. As UV-crosslinkable PSAs have a pendant photoinitiator group in the PSA polymer backbone, crosslinking occurs with neighboring groups when they are UV-irradiated. If the copolymerized photoinitiator is type II, such as a benzophenone, it abstracts hydrogen from the hydrocarbon when exposed to UV through the crosslinking mechanism (Scheme 17.2) [7]. If the hydrocarbon contains hydroxyl, amine, and ether groups, the benzophenone group can abstract more hydrogen from these groups [10]. The objective of this study was to prepare a UV-crosslinkable acrylic PSA that is able to cure at low UV doses and to elucidate the effect of crosslinking on PSA performances at varying UV doses. The PSAs used in this study were synthesized by solution polymerization method using 2-ethylhexyl acrylate (2-

251

252


Scheme 17.2 UV-crosslinking reaction of PSA via a hydrogen abstraction process (R' = alkyl, alkylene, hydroxyl, or amine group) [7].

EHA), vinyl acetate (VAc) and acrylic acid (AA), and with varying contents of 2hydroxyethyl methacrylate (2-HEMA) and an unsaturated benzophenone derivative (P-36). The UV-crosslinking behavior of the PSAs was monitored by FTIRATR and PSA performance was evaluated by probe tack, peel strength, and shear adhesion failure temperature (SAFT) after exposure to various UV doses.

17.2 Materials and Methods 17.2.1 Preparation of UV-Crosslinkable Acrylic PSA

The 2-ethylhexyl acrylate, vinyl acetate, acrylic acid, 2-hydroxyethyl methacrylate, and 2,2'-azobisisobutyronitrile (AIBN, as initiator for polymerization) used for preparing UV-crosslinkable acrylic PSAs were obtained from Junsei Chemicals Co. and used as received. The photoinitiator P-36 containing a double bond was obtained from SK UCB Co., Korea. The typical synthetic method was as follows (see Table 17.1). 2-EHA (120 g), VAc (22.5 g), AA (7.5 g), AIBN (0.3 g), and ethyl acetate (75 g) were all mixed in a 500 mL four-necked flask equipped with a stirrer, dropping funnel, and thermometer. The polymerization reaction was initiated at 70 8C. After this temperature had been maintained for 1 h, ethyl acetate (75 g) containing P-36 (1.5 g) was gradually added to the flask while it was stirred for 1 h, and then the polymerization was carried out at 70–75 8C for another 5 h. All of the syntheses were carried out in a flask wrapped in aluminum foil

17.2 Materials and Methods Table 17.1 PSA formulations synthesized in this work. a)

Code

2-EHA [wt%]

VAc [wt%]

AA [wt%]

2-HEMA [phr] b)

P-36 [phr]

SH0P1 SH3P1 SH6P1 SH9P1 SH3P05 SH3P2

80 80 80 80 80 80

15 15 15 15 15 15

5 5 5 5 5 5

0 3 6 9 3 3

1 1 1 1 0.5 2

a) b)

The solids content of all the synthesized PSAs was 50 wt%. phr: parts per hundred resin, the phr unit can be converted into wt% by the following equation: wt% = (100 ´ phr)/(100+phr).

to minimize the light effect. The final product (SH0P1) was a colorless, viscous liquid and could be used as a UV-crosslinkable acrylic PSA. The other compositions of the UV-crosslinkable acrylic PSAs synthesized are listed in Table 17.1. 17.2.2 Preparation of PSA Samples and UV Curing

All UV-crosslinkable acrylic PSAs were coated on a poly(ethylene terephthalate) (PET) (SK Chemical Co. Ltd.) film 25 lm thick using a No. 26 K-bar, kept at room temperature for 1 h and then dried in an oven at 90 8C for 2 h. These dried PSA films were kept at 22 ± 2 8C and 60 ± 5% RH for 24 h before testing and were cured in a conveyor belt type of UV curing machine equipped with a high-pressure mercury lamp (100 W cm–1, power loaded, main wavelength 340 nm). In order to avoid heating effects from the mercury lamp, a cold mirror was used as a reflector. The UV doses (0, 210, 630, 1050, 1470, 1890, 2310, and 2730 mJ cm–2) were measured with an IL 390C Light Bug UV radiometer (International Light Inc.). 17.2.3 FTIR-ATR Spectroscopy

The IR spectra were obtained using a Nicolet Magna 550 Series II FT-IR (Midac Co., USA) equipped with 458 zinc selenide (ZnSe, n = 2.4) attenuated total reflectance (ATR). Spectra were collected for 32 scans at a resolution of 16 cm–1 between 650 and 4000 cm–1. The spectrometer was linked to a PC equipped with Omnic E.S.P. 5.2 software to allow the automated collection of IR spectra and to integrate the peak area.

253

254


17.2.4 DSC Measurement

Differential scanning calorimeter (DSC, TA Q-1000) measurements were carried out using a constant sample weight of 20 mg, in aluminum pans under a flowing atmosphere of nitrogen (50 cm3 min–1) at a heating rate of 10 K min–1 from –70 to 150 8C. 17.2.5 PSA Performance

PSA performance was evaluated by probe tack, peel strength, and SAFT. The probe tack and peel strength tests were conducted using a Texture Analyzer (Stable Micro Systems, TA-XT2i). The probe tack test with a type 304 stainless steel cylinder probe of diameter 5 mm was carried out at a separation rate of 0.5 mm s–1 under a constant pressure of 980 N m–2 and a dwell time of 1 s on each PSA layer. The surface roughness height of the probe was 250 nm. The peel strength test was performed after applying the PSA sample (25 mm ´ 300 mm: width ´ length) on type 304 stainless steel (50 ´ 150 mm and 1.5 mm thick) of which the surface roughness height was 40 ± 5 nm, and keeping the sample at room temperature for 24 h at an angle of 1808 with a cross-head speed of 300 mm min–1. For SAFT, the 25 mm ´ 150 mm PSA samples were pressed onto stainless steel which was same as that used for the peel strength test (bonding area: 25 mm ´ 25 mm) by two passes of a 2 kg rubber roller at a rate of 10 ± 0.5 mm s–1. The stainless steel was mounted vertically and a load of 1 kg was attached to the free end of the PSA sample, while the bonded area was adhered to the stainless steel substrate. The failure temperature was measured as the temperature was increased at 0.4 K min–1.

17.3 Results and Discussion 17.3.1 FTIR-ATR Measurements

Benzophenone, type II, is the most widely used photoinitiator in photochemistry. When exposed to UV, the carbonyl group in benzophenone is excited to a triplet state of n, p* configuration through a singlet state. These triplets abstract hydrogen from suitable hydrogen-donating groups such as alcohol, alkylbenzene, hydrocarbon, and tributyl hydride, as shown in Scheme 17.1 [10, 11]. Therefore, if a benzophenone derivative containing a vinyl double bond is copolymerized via radical polymerization in the polymer backbone, it can work as a crosslinker when exposed to UV (Scheme 17.3) [12]. Because the conversion of the C=O group in benzophenone leads to the loss of conjugation between the carbonyl group and the aromatic ring upon UV ex-

Scheme 17.3 Photochemical crosslinking of an acrylic PSA via the triplet state of copolymerized benzophenone groups in the PSA polymer backbone [12].

17.3 Results and Discussion 255

256


Fig. 17.2 IR spectra of the copolymerized benzophenone groups in an acrylic PSA: UV dose at (a) 0 mJ cm–2; (b) 2730 mJ cm–2.

posure, the kinetic profile of the UV-crosslinkable acrylic PSAs synthesized containing a benzophenone group can be evaluated by observing their characteristic absorption band at about 1600 cm–1 (C=C in the C6H5 of the benzophenone group) with increasing UV dose [12–14]. Fig. 17.2 shows the FTIR-ATR spectra of synthesized PSA at UV doses of 0 and 2730 mJ cm–2. The peak of the C=C in the aromatic ring almost vanished at a dose of 2730 mJ cm–2. The loss of conjugation between the C=C and C=O in benzophenone with increasing UV dosage is caused by the rising concentration of triplet states and by the abstraction of hydrogen from neighboring hydrogen donors such as the hydroxyl groups; as in 2-HEMA, the rate of the loss of conjugation between C=C and C=O will increase. Hence, an increased content of both photoinitiator and 2-HEMA gives rise to the loss of conjugation between C=C and C=O in benzophenone. In order to evaluate these two effects on the kinetic profile of UV-crosslinkable acrylic PSA, the relative concentration of C=C in the benzophenone was calculated by integrating the peak area between 1577 and 1604 cm–1. Fig. 17.3 shows the decrease in the relative contents of the C=C in the benzophenone of SH0P1 and SH3P1 with increasing UV dose. The relative content of the C=C in benzophenone of SH3P1 decreased faster than that of SH0P1 at 210 mJ cm–2 (Fig. 17.3). However, at UV doses above 630 mJ cm–2, there was no difference in the relative contents of the C=C in the benzophenone of SH3P1 and SH0P1. These effects are due to the higher hydrogen-donating ability of 2HEMA, which has a hydroxyl group. In Fig. 17.4, the influence of the composition of the benzophenone is shown. The relative contents of the C=C in the benzophenone of SH3P2 were lower than those for SH3P1 at a UV dose of 210 mJ cm–2. This effect may be due to the higher content of benzophenone in PSA, which increased the efficiency of


Fig. 17.3 Change of relative concentration of C=C bonds in SH0P1 and SH3P1 with variation in UV dose.

Fig. 17.4 Change of relative concentration of C=C bond in SH3P1 and SH3P2 with variation in UV dose.

UV curing at low UV doses. But at UV doses above 630 mJ cm–2, the relative C=C contents were almost the same. This may be due to the increased crosslinking density, which reduced the mobility in the PSAs. This phenomenon confirmed the result of many previous studies, that the maximum UV-curing rate was shifted to an earlier stage of UV irradiation when the photoinitiator content was increased. After this point, however, the increased crosslinking density reduced the molecular mobility and so the curing rate decreased [15–17].

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17.3.2 PSA Performance 17.3.2.1 Probe Tack Fig. 17.5 shows the change of probe tack with varying 2-HEMA contents at various UV doses. The probe tack of SH3P1, SH6P1, and SH9P1 decreased dramatically at a UV dose of 210 mJ cm–2, but at above 630 mJ cm–2 the probe tack decreased only slightly. The relative decrease (D) in probe tack at 210 mJ cm–2 compared with probe tack before UV exposure was calculated by means of Eq. (1).

D% 1

tack210mJ=cm2 tack0mJ=cm2

100

1

The decreases for SH0P1, SH3P1, SH6P1, and SH9P1 were 0.6%, 7.3%, 8.4%, and 11.8%, respectively. Before UV irradiation, as probe tack values generally depend on the PSA composition, the greater the 2-HEMA contents, the stronger was the polarity, which caused both probe tack and peel strength to increase. The later values were calculated to compare the relative change in probe tack values. These results can be explained by the fact that the increased 2-HEMA acted as a good hydrogen donor (more than any other hydrocarbon in the UV-crosslinking system); the crosslinking density was increased at low UV doses, thereby reducing the molecular mobility and decreasing the probe tack. The effects of photoinitiator concentration on the probe tack are shown in Fig. 17.6. As the rate of crosslinking of SH3P2, which had a higher content of photoinitiator, was faster than that of SH3P05 at 210 mJ cm–2 as mentioned above, the in-

Fig. 17.5 Change of probe tack in SH0P1, SH3P1, SH6P1, and SH9P1 with variation in UV dose.


Fig. 17.6 Change of probe tack in SH3P05, SH3P1, and SH3P2 with variation in UV dose.

creased crosslinking density prevented the PSA from wetting the probe, which consequently reduced the probe tack more in SH3P2. The decreased values for SH3P2 and SH3P05 were 15.7% and 1.2%, respectively. In order to perform appropriately as a PSA, the PSA polymer must have a sufficiently low Tg; depending on its application, it may be far below room temperature [3]. The Tg of a PSA can be adjusted by formulating low-Tg material (soft polymer: elastomer) and high-Tg additives (hard polymer or oligomer). As the Tg of the PSA is lowered, tack and elongation at break increase but hardness and tensile strength decrease [18]. Table 17.2 shows the Tg data of each PSA sample for various UV doses. Before UV irradiation, the Tg of the PSA formulations was about –60 8C. With increasing UV dose, the Tg rose for each PSA. These results correspond to the probe tack results, in that before UV irradiation (lower Tg) PSA samples showed a high probe tack, and after UV irradiation (higher Tg) the probe tack was reduced for all the samples. With increasing UV

Table 17.2 Variation in glass transition temperature Tg [8C] of PSAs with UV dose.

UV dose [mJ cm–2]

SH0P1 SH3P1 SH6P1 SH9P1 SH3P05 SH3P2

0

210

1470

2310

–61.1 –61.0 –62.3 –61.9 –62.1 –61.3

–45.6 –45.4 –44.5 –49.6 –48.7 –40.0

–42.7 –42.6 –42.6 –47.8 –46.3 –40.8

–42.2 –42.8 –43.6 –46.0 –46.6 –40.3

259

260


dose, however, the increment in Tg became negligible and the tack values did not vary from sample to sample at higher UV doses. These Tg results and probe tack results show similar tendencies.

17.3.2.2 Peel Strength The peel strength results of SH0P1, SH3P1, SH6P1, and SH9P1 are shown in Fig. 17.7. At UV doses above 630 mJ cm–2, all PSAs gave equal values, but on exposure to low UV doses (210 mJ cm–2) they showed different behaviors in peel strength. The peel strength values of SH6P1 and SH9P1, which had a higher content of 2-HEMA, were dramatically reduced by 53.3% and 52.7%, respectively, at 210 mJ cm–2. Contrarily, the peel strength values of SH0P1 and SH3P1 were reduced only by 8.9% and 15.1%, respectively, at the same UV dose. These results are fully comparable with those for the probe tack, in which the increased contents of 2-HEMA caused the PSA crosslinking in the early stages of UV exposure to produce a rapid decrease in peel strength. For example, the high crosslinking density of PSAs further increases the volume contraction and the storage modulus, thereby decreasing the peel strength to low values [19, 20]. The effects of the photoinitiator content in PSAs are shown in Fig. 17.8. As expected, the higher photoinitiator content in SH3P2 produced a rapid decrease in peel strength by 69.8% at a UV dose of 210 mJ cm–2, whereas SH3P05 gave a 9.1% reduction in peel strength at the same dose. Probe tack and peel strength test methods are almost same, except for test area and test time, in that the measurements are taken with the PSA removed from the substrate. So the peel strength is influenced by the Tg also, so that if the Tg of a PSA is sufficiently low, the wettability will increase after applying it

Fig. 17.7 Change of peel strength in SH0P1, SH3P1, SH6P1, and SH9P1 with variation in UV dose.


Fig. 17.8 Change of peel strength in SH3P05, SH3P1, and SH3P2 with variation in UV dose.

to a substrate (stainless steel, PVC, PE, PP, etc.) and as a result the peel strength will increase [19]. The Tg data shown in Table 17.2 and the peel strength results showed a similar tendency. Before UV irradiation of the PSA (lower Tg) all the PSAs exhibited higher peel strengths, and after UV irradiation (higher Tg) the peel strengths were lowered. As discussed in Section 17.3.2.1, at higher UV doses for the PSAs, the increment in Tg was negligible; this also influenced the peel strength so there the peel strength did not change after UV irradiation.

17.3.2.3 Shear Adhesion Failure Temperature (SAFT) In general, an acrylate copolymer which has no crosslinking, or which is crosslinked only by hydrogen bonding, does not show an adequate PSA performance at elevated temperature, so that chemical crosslinking is required to provide high shear strength [21]. The shear strength of PSA is measured by a dynamic or static method [22]. In the dynamic shear test, the failure temperature, SAFT, is measured by pulling the PSA downward under a constant force from a vertically placed test substrate in a direction parallel to the surface of bonding, while the temperature is gradually raised. On the other hand, in the static method a time is evaluated using a constant load at a fixed temperature [22, 23]. The SAFT results of PSAs with varying contents of 2-HEMA are shown in Fig. 17.9. With an increasing UV dose, the PSAs showed higher SAFT values. SH9P1, especially, with higher 2-HEMA contents, showed high SAFT values at low UV doses and reached a SAFT level of 150 8C earlier than SH3P1 and SH0P1. This means that 2-HEMA acted as a good hydrogen donor to the benzophenone group, thereby allowing the increased crosslinking density at the lower UV doses to increase the cohesion of the PSA polymer. Fig. 17.10 presents the

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262


Fig. 17.9 Change of SAFT in SH0P1, SH3P1, and SH9P1 with variation in UV dose.

Fig. 17.10 Change of SAFT in SH3P1 and SH3P2 with variation in UV dose.

effects of photoinitiator concentration on SAFT at various UV doses. As the rate of photopolymerization is proportional to the photoinitiator concentration, the increased amount of photoinitiator in the PSA produced a high crosslinking density at low UV doses. As a result, the SAFT values of SH3P2 were higher than those of SH3P1.

References

17.4 Conclusions

Of the various methods used to manufacture UV-curable PSA, that of preparing a UV-crosslinkable acrylic PSA was adapted to evaluate the UV-curing behavior and adhesion performance at various UV doses. The UV-crosslinkable PSAs were copolymerized with acrylate monomers and an unsaturated photoinitiator, P-36, with varying contents of 2-HEMA and photoinitiator. From the FTIR-ATR results, the rate of crosslinking of the PSAs which had higher contents of 2-HEMA and photoinitiator was greater than that of the PSAs with low contents of 2-HEMA and photoinitiator. The PSA performance, probe tack, peel strength, and SAFT were also affected by the curing rate. With increasing concentrations of 2-HEMA and photoinitiator in the PSAs, the probe tack and peel strength were rapidly reduced at the beginning of UV exposure, whereas the SAFT values were increased.

References 1 D. Satas (Ed.), Handbook of Pressure Sen-

2 3

4 5 6 7 8 9 10

11

12

sitive Adhesive Technology and Applications, Satas & Associates, Warwick, 2002, Chapter 4. H. Miller, Adhesives Sealants Ind. 2002, May, 34–41. D. Satas (Ed.), Handbook of Pressure Sensitive Adhesive Technology and Applications, Satas & Associates, Warwick, 2002, Chapter 19. Z. Czech, Polym. Int. 2003, 52, 347–357. H. S. Do, S. E. Kim, H. J. Kim, J. Adhes. Sci. Technol. submitted. Z. Czech, R. Milker, J. Appl. Polym. Sci. 2003, 87, 182–191. Z. Czech, Europ. Polym. J. 2004, 40, 2221–2227. Z. Czech, Int. J. Adhes. Adhesives 2004, 24(2), 119–125. H. S. Do, S. E. Kim, H. J. Kim, J. Adhes. Interface, in press. J. F. Rabek, Mechanisms of Photophysical Processes and Photochemical Reactions in Polymers, John Wiley & Sons, Chichester, 1987, Chapters 7 and 11. C. G. Roffey, Photogeneration of Reactive Species for UV Curing, John Wiley & Sons, Chichester, 1997, Chapter 5. T. Scherzer, A. Tauber, R. Mehnert, Vibrational Spectroscopy 2002, 29, 125–131.

13 N. B. Colthup, L. H. Daly, S. E. Wiberley,

14

15 16 17 18

19 20 21 22

23

Introduction to Infrared and Raman Spectroscopy, Academic Press, San Diego, 1990, p. 296. G. Socrates, Infrared and Raman Characteristic Group Frequencies, John Wiley & Sons, Chichester, 2001, p. 161. T. Scherzer, U. Decker, Vibrational Spectroscopy 1999, 19, 385–398. L. Lecamp, B. Youssef, C. Bunel, P. Lebaudy, Polymer 1997, 38, 6089–6096. H. S. Do, H. J. Kim, Y. K. Lee, Adhes. Interface 2003, 4(3), 14–20. I. Benedeck, L. J. Heymans, Pressure Sensitive Adhesives Technology, Marcel Decker, New York, 1997, Chapter 3. K. Ebe, H. Seno, K. Horigome, J. Appl. Polym. Sci. 2003, 90, 436–441. H. S. Do, S. E. Kim, H. J. Kim, Adhes. Interface 2004, 5(4), 1–7. Z. Czech, Int. J. Adhes. Adhesives 2004, 24, 503–511. I. Benedeck, L. J. Heymans, Pressure Sensitive Adhesives Technology, Marcel Decker, New York, 1997, Chapter 6. N. Nakajima, R. Babrowicz, E. R. Harrell, J. Appl. Polym. Sci. 1992, 44, 1437–1456.

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18 Contribution of Chemical Interactions to the Adhesion Between Evaporated Metals and Functional Groups of Different Types at Polymer Surfaces J. Friedrich, R. Mix, and G. Kühn

Abstract

Single-type functionalizations with different types of functional groups at polypropylene (PP) and polytetrafluoroethylene (PTFE) surfaces were achieved using, instead of a simple plasma modification, either a combined plasmachemical–chemical process or the pulsed plasma-initiated homo- or copolymerization of monomers carrying functional groups. The combined process consists of O2 plasma pretreatment and wet-chemical reduction of O functional groups to OH groups using diborane, vitride (sodium bis(2-methoxyethoxy)aluminum hydride), or LiAlH4. The high degree of retained chemical structure and functional groups during the low-power pulsed plasma polymerization was found to be a prerequisite for producing well-defined, adhesion-promoting plasma polymer layers as model surfaces with high concentrations of exclusively or predominantly one type of functional group, such as OH, NH2, or COOH. The maximum concentrations of functional groups were found to be 31 OH, 21 NH2 or 25 COOH groups/100 C atoms using allyl alcohol, allylamine, or acrylic acid, respectively, as monomers in the plasma polymerization process and 14 OH groups/100 C atoms by applying the combined O2 plasma/diborane reduction process. To vary the density of functional groups, a so-called plasma-initiated gas-phase radical copolymerization with ethylene or styrene as a “chain-extending” comonomer, or butadiene as “chemical crosslinker” was employed. The peel strength of evaporated aluminum layers on unspecifically oxygen-plasma functionalized polypropylene (PP) and polyethylene (PE) shows in each case a maximum at 20 O per 100 C atoms. Initially the peel strength increased linearly with the concentration of functional groups when PP or polytetrafluoroethylene (PTFE) substrates were coated with plasma polymers or copolymers carrying a single type of adhesion-promoting functional groups. The ranking of the adhesion-promoting effect is CH2 < NH2 OH < COOH, and corresponds to the tendency to form chemical bonds between aluminum and the different functional groups.


266

18 Contribution of Chemical Interactions to the Adhesion Between Metals and Functional Groups

18.1 Introduction 18.1.1 Interactions Between Metal Atoms and Functional Groups at Polymer Surfaces

The adhesion of metal layers deposited onto polymer surfaces is determined by the concentration and the bond strength of the chemical and physical interactions between the metal atoms and the functional (polar) groups at the polymer surfaces. Each type of functional group produces individual metal–polymer interactions, and makes a specific contribution depending on its concentration to the interfacial adhesion and consequently to the related shear or peel strength of metal–polymer systems (see Fig. 18.1). Thus for each type x of metal–functional group interaction rx, the work of adhesion Wax is calculated with Eq. (1), with A = area, L = Loschmidt’s constant, and ci = concentration. Wax A L ci rm

1

The total work of adhesion Wa follows by summation over all types of interaction [Eq. (2)], presenting the correlation with the macroscopically measured adhesive bond strength, here the peel strength. Wa A L

m X

ci ri

2

i1

The aim of this work was to prepare a well-defined system for investigating the individual contribution of each type of metal–functional group interaction by designing polymer surfaces with exclusively one type of functional group. Then, the interfacial adhesion is related to only that type of metal–functional

Fig. 18.1 Unspecific and specific functionalization before and after interactions with metal layers.

18.1 Introduction

group interaction and thus it is depending exclusively on the respective functional group concentration as seen above. However, additional unspecific contributions to the adhesion strength due to van der Waals interactions, hydrogen bonding, and interdiffusion processes must also be considered [1]. 18.1.2 Preparation of the Plasma-Modified Polymer Surfaces

Four ways of polymer surface functionalization were considered and the following surface compositions were expected (see Fig. 18.2): 1. unspecific functionalization in the low-pressure oxygen plasma leading to different types of O-functional groups at various concentrations 2. improvement of this unspecific O-functionalization to a specific OH-functionalization using a post-plasma chemical reduction of carbonyl-containing groups to OH groups 3. a pulsed plasma polymerization technique that provides homopolymer layers carrying adhesion-promoting functional groups 4. pulsed plasma polymerization of copolymers from monomers with functional groups and “neutral” monomers for “chain extension”; in this way, the concentration of single-type functional groups can be adjusted continuously. The unspecific functionalization of polymer surfaces (1) was first described for polyolefin surfaces in 1956 by Rossmann [2]. X-ray photoelectron spectroscopy (XPS) studies especially have cast much light on the nature of this unspecific oxygen-plasma modification of polymers [3–6]. It was found that the low-pressure oxygen-plasma modification rapidly changes from pure formation of functional groups at the polymer surface to oxidative degradation, crosslinking, and etching of near-surface layers of the polymer when the treatment exceeds ca. 2 s [7]. For adhesion promotion, the nature of the functional groups introduced was not considered to be critical as long as all of them possess sufficient polarity and thus produce physical interactions between metal atoms and polar groups

(1)

(2)

(3)

(4)

Fig. 18.2 Functionalized polymer surfaces prepared by applying the processes 1–4 listed in the text.

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[8]. Aside from the numerous possible interactions between metal atoms and the different types of O-functional groups, a few investigators report on the formation of chemical bonds, most often of the Me–O–C or Me–C types [9–11]. New possibilities for the investigation by method (2) opened when the mixture of functional groups was unified after O-plasma treatment by the reduction of carbonyl groups to OH groups and by the simultaneous hydroboration of double bonds to OH groups using diborane (B2H6)/H2O2 [Eqs. (A, B)] [12]. Such a reduction/hydroboration of different functional groups can be performed with a maximum yield of 14 OH per 100 C atoms. The conversion to OH groups was performed as a wet-chemical reduction with B2H6, vitride (sodium bis(2-methoxyethoxy)aluminum hydride), or LiAlH4 [13]. 2C=O B2 H6 ! 2 CH OBH2 H2 SO4 ! 2 CH OH

A

2C=C B2 H6 ! 2 H C C BH2 2 H2 O2 ! 2 H C C OH

B

Route (3) to improve the adhesion properties of polymers was opened by the deposition of a thin adhesion-promoting plasma polymer layer using the pulsed plasma technique. The pulsed plasma polymerization is preferred because it enhances the pure chemical chain propagation. In contrast, the “continuous-wave plasma polymerization,” first used by Goodman [14] and interpreted by Yasuda as “atomic polymerization” [15], is dominated by monomer fragmentation followed by random fragment recombination polymerization. Thus, irregular structured polymers were produced. Using the new pulse technique [16–20] the plasma polymers were formed predominantly by chemical chain propagation, resulting in chemically better defined polymer structures. Theoretically, one short plasma pulse should be enough to start a pure chemical radical chain polymerization in the gas/plasma phase or a graft polymerization of suitable monomers, such as acrylic, vinyl, or diene monomers, at the substrate surface. However, the number of monomers sticking at radical sites of the growing chains is lower by the factor 108 if low-pressure plasmas (1–20 Pa) are compared with chemical polymerizations in the liquid phase [21]. The radical chain reaction is frequently terminated by disproportionation, a radical–radical recombination, or a reaction with oxygen molecules from the rest of the gas. Therefore, pulsed plasmas with long plasma-off periods and short plasma pulses offer a good compromise to produce chemically formed polymer structures with a minimum of irregularities that are produced during the “plasmaon” periods. In the plasma-produced homopolymers from allyl alcohol, the first tests had shown that 95% of the theoretically expected OH groups were retained. Only a small concentration of C=O “impurities” could be verified [22–24]. The detection of primary amino groups in allylamine homopolymers is difficult because the derivatization reaction needed for quantification is incomplete [25]. Moreover, the functionalization of surfaces with NH2 groups is superimposed by the attachment of O-functional groups because of the extensive oxidation of the car-

18.1 Introduction

bon atom in the a-position to the amino group when exposed to air [26]. Homopolymeric layers made from acrylic acid are stable on exposure to air. However, fragmentation of the acid group during the plasma process limits the yield of COOH groups. 18.1.3 Interactions Between Evaporated Al and Functional Groups

Four different functional groups, CH2, NH2, OH, and COOH, on polypropylene (PP) and on polytetrafluoroethylene (PTFE) (process 3) as well as the unspecific O-functionalization of these polymers with different types of O-functional groups (process 1) were tested for their adhesion-promoting properties toward aluminum top coatings. Aliphatic groups do not form significant interactions with Al [27, 28] but under appropriate conditions Al–C formation was also observed [29–32]. The existence of chemical interactions between Al and primary amino groups is also controversial. An amino group forms even at low temperature an Al ?NH2 interaction (r-complex) which may rearrange to Al–N [33] or with Mg to Mg–N [34]. The global improvement of Al adhesion by introduction of N-containing groups was demonstrated by Andre et al. [35]. The OH group in poly(vinyl alcohol) (PVA) reacts with trimethylaluminum ((CH3)3Al) and also with evaporated metallic Al forming Al–O–C complexes [36, 37], as with evaporated Ag [38]. Other authors interpret the electron transfer from Al to different functional groups present at a polymer surface as acid–base interactions [39]. It was also shown that a few atom layers of deposited Al (ebeam, thermal evaporation) were oxidized to Al2O3 by simultaneous reduction of all the O-functional groups present at O2-plasma modified polypropylene surfaces [40, 41]. This redox mechanism was confirmed by Silvain et al. [42] and Ding et al. [43], who found that Al reacts with the F of a fluoropolymer and forms AlF3 at the interface. Al adheres very strongly to COOH-group modified polymers [44, 45]. The resulting chemical bond may be Al+COO– (salt-like) [37]. Another mechanism prefers the interaction of Al with the carbonyl site first and, more slowly, with the C–O site [46], as was also found with COOR groups [32, 47, 48]. A third mechanism is the redox reaction of Al and carboxylic groups, thus forming Al2O3 and destroying the carboxylic groups [40]. 18.1.4 Adhesive Bond Strength and Concentration of Functional Groups

Low concentrations of appropriate reactive functional groups (0.001–0.1 groups per polymer repeat unit), introduced by a chemical copolymerization onto the polymer surface, are found to increase the adhesion strength with a negligible influence on the bulk properties [49]. The type of functional group determines the resulting specific interactions with metals [50]. The concentration of func-

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tional groups at the polymer surface should influence which mode of metal film formation, i.e., cluster formation (Vollmer–Weber mode) or homogeneous film formation (Frank van der Merve mode), is present. This aspect has been neglected and must be considered in future work. However, it can be expected that at low concentrations of polar functional groups at the surface island growth is dominant, and at high concentrations of polar groups the spreading of metal atoms is improved. To vary the density of functional groups, a radical copolymerization with ethylene as a “chain-extending” comonomer, or butadiene as a “chemical crosslinker,” was performed using the pulsed plasma. As in classic copolymerization, the main problem was to find compatible pairs of copolymers, to realize copolymerization in a broad range of mixing ratios of comonomers, and to avoid the dominance of homopolymerization of an individual comonomer [51, 52]. It was pointed out that this process was a “pulsed plasma-initiated radical gas-phase copolymerization,” which strongly contrasted with the simple, so-called “plasma copolymerization”. That “old” process of plasma copolymerization also allows for the reaction of nonpolymerizable (chemically inert) “monomers,” as previously demonstrated [53–55]. The most significant disadvantage of the “old” continuous-wave (cw) plasma copolymerization process is the almost complete monomer fragmentation to atoms and small fragments due to the amount of energy consumed per monomer during residence in the plasma zone, which is generally sufficient to break all the chemical bonds in the monomer molecule [56–58]. The fragmentation is followed by random polyrecombination of fragments and atoms to an irregular, undefined structure. Additional defects in the deposited plasma polymer are produced by permanent exposure of the growing copolymer layer to the plasma UV radiation during the deposition process [59]. The “new” pulsed plasma copolymerization process minimizes these disadvantages and increases significantly the fraction of copolymers that are produced in a purely chemical way during the “plasma-off” period between two plasma pulses. However as was pointed out earlier, with the “new” method also, irregularities were produced during the plasma pulses but in much smaller quantities [52].


Polypropylene (PP) films (100 lm from Goodfellow, UK, or 300 lm from CibaGeigy, Switzerland) as well as polytetrafluoroethylene (PTFE) foil (1 mm from DuPont, USA) were ultrasonically cleaned in a diethyl ether bath for 15 min. The monomers used were: acrylic acid, allyl alcohol, and allylamine (all > 99%


from Merck KGaA, Darmstadt, Germany). These monomers were distilled before use. The other monomers, ethylene and 1,3-butadiene (Messer-Griesheim GmbH, Germany) were used as received. 18.2.2 Plasma Pretreatment of Polymers

Surface functionalization of PP films in the O2 plasma was performed in the cw mode. PP films which were coated with plasma polymer layers of functional-group carrying monomers had been used without any additional plasma pretreatment. PTFE films were exposed first to H2 radio-frequency (RF) plasma (cw mode) for 1–1800 s at pressure p = 6 Pa and power P = 300 W, followed by the deposition of adhesion-promoting plasma polymer layers. 18.2.3 Deposition of Adhesion-Promoting Plasma Polymer Layers

Plasma polymer layers were deposited in a cylindrical plasma reactor of 50 dm3 volume. The design of the plasma reactor has been described in detail in Ref. [60]. The reactor was equipped with a pulsed RF (13.56 MHz) generator with an automatic matching unit and an RF bar antenna (length = 35 cm). The substrate foil was mounted on a continuously rotating, grounded steel cylinder (Æ = 10 cm, length = 35 cm, rotating frequency = 0.5–1 s–1) at 10 cm from the RFpowered electrode. The duty cycle of pulsing was adjusted to 0.1 and the pulse frequency to 103 Hz. The power input was varied between 100 and 300 W. Mass flow controllers for gases and vapors, a heated gas/vapor distribution system in the chamber, and control of pressure and monomer flow by varying the speed of the turbomolecular pump were used. The gas flow was adjusted to 75– 125 sccm and the pressure was kept at 26 Pa. The deposition rate was measured by a quartz microbalance. 18.2.4 Surface Analysis

The XPS and IR analyses have been described in detail elsewhere [61]. Here, only some important facts are summarized. The XPS data acquisition was performed with a Sage 150 Spectrometer (Specs, Berlin, Germany) using nonmonochromatized MgKa or AlKa radiation with 12.5 kV and 250 W settings at a pressure of &10–7 Pa in the analysis chamber. XPS spectra were acquired in the constant analyzer energy (CAE) mode at 908 take-off angle. Peak analysis was performed using the peak fit routine from Specs. The FTIR spectra were recorded with a Nexus instrument (Nicolet, USA) using the ATR (attenuated total reflectance, 458 angle of incidence) technique with a diamond or a Ge cell (“Golden Gate,” Specac, Kent, UK). The IR signal comes from a near-surface layer of the polymer film. The information depth de-

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pends on the material used as the ATR crystal and amounts to about < 1.5 lm using germanium and a 2.5 lm diamond applicator . 18.2.5 Labeling of Functional Groups

For exact XPS quantification of the functional groups, derivatizations were performed using trifluoroacetic anhydride (TFAA) [62, 63] or m-trifluoromethylphenyl isocyanate (TMPI) for the OH groups, pentafluorobenzaldehyde (PFBA) or 4-trifluoromethylbenzaldehyde (TFMBA) for the NH2 groups [64, 65], and trifluoroethanol (TFE) in the presence of dicyclohexylcarbodiimide for the COOH groups [66, 67]. The number of functional groups was calculated by considering the percentage of fluorine introduced (the F1s peak intensity in XPS) and the theoretical stoichiometry of the derivatized polymer. It was assumed that the outermost layer (&3 nm), analyzed by XPS, was uniformly derivatized. The completeness of the derivatization and the absence of unconsumed functional groups inside the deposited polymer layer (&50 nm) were checked by measuring ATR spectra of the surface of plasma polymers and inspecting the respective bands. The same spectra were measured when the plasma polymers were deposited directly onto ATR crystals and then derivatized. Moreover, C1s peak fitting (CF3, COOR) was achieved using the peak fit routine from Specs. 18.2.6 Contact Angle Measurements

Static contact angle measurements were performed on sessile drops of water, formamide, ethylene glycol, benzyl alcohol, and diiodomethane as test liquids. A contact-angle measuring system G2 (Krüss, Hamburg, Germany) including the appropriate software was used. The software employed for the calculation of surface energy is based on the approaches of Owens and Wendt as well as Rabel and Kaelble [68–70]. The surfaces of plasma polymer layers were inspected by atomic force microscopy (AFM) to obtain information on their roughness. It was found that the layers were very flat on the micrometer scale (RRMS = 1–2 nm) for a thickness of £100 nm. 18.2.7 Metal Deposition

To measure the Al–polymer (PTFE, PP) peel strength, the plasma polymerization was performed using a plasma reactor (Ilmplasma 1200; Saskia, Ilmenau, Germany). Then, the plasma polymer coated polymer samples were transferred into a separate electron beam metallizer (Auto 306; Edwards, UK). The thickness of deposited aluminum layers was adjusted to 150 nm using a quartz microbalance.

18.3 Results

18.2.8 Peel Strength Measurements

The metal peeling technique for Al–PP and Al–PTFE systems followed Dupont’s preparation and peeling procedure. The Al side of the composite was fixed onto double-faced Scotch® tape, which was pressed onto a steel plate [71–73]. Then, a 908 peel test was carried out at a peel speed of 25 mm min–1 for all the Al– polymer systems. The standard deviation of the peel strength varied between 10 and 15%. Low peel strength was linked with interfacial failure, moderate values with partial Al peeling, and high peel strength with complete retention of Al on the polymer side. In such a case, interfacial failure at the interface between an adhesion-promoting layer and the polymer substrate, within the bulk of the adhesion-promoting plasma polymer layer, or inside the polymer substrate was also observed. The thickness of each of the adhesion-promoting pulsed plasma polymer layers was adjusted to 150 nm. After peeling of Al–polymer systems with adhesion-promoting plasma copolymer layers, the peeled surfaces were inspected again with XPS to determine the locus of failure, i.e., whether the peel front propagated along the interface (interfacial failure) or within the material (cohesive failure).

18.3 Results 18.3.1 Production of Polymer Surfaces Containing Functional Groups

By means of OH- and COOH-containing plasma polymer layers the qualification of these layers as models of single-type functionalized adhesion promoters with variable concentrations of functional groups should be proved. The plasma-initiated copolymerization of acrylic acid with ethylene or 1,3-butadiene is shown in terms of measured COOH concentration as a function of the composition of the comonomer mixture in Fig. 18.3. Depending on the co-monomer reactivity, a more linear correlation (butadiene), or a parabolic behavior (ethylene), between precursor composition and COOH groups produced was observed. For each type and concentration of functional group, its concentration was determined by chemical derivatization followed by XPS analysis as described in Section 18.2.5. As shown in Fig. 18.4, similar tendencies were found for the allyl alcohol copolymerization with ethylene, butadiene, or styrene. Here, the curves are observed to progress from a parabolic (ethylene) to a nearly linear correlation (butadiene) and to an anti-parabolic behavior (styrene) between measured OH group concentrations and the stoichiometry of the precursor mixture. By considering these curves it was possible to adjust to a defined concentration of functional groups in the deposited layer. Allylamine copolymer and pure

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18 Contribution of Chemical Interactions to the Adhesion Between Metals and Functional Groups Fig. 18.3 Yield in COOH groups as a function of the composition of the precursor mixture.

Fig. 18.4 Dependence of yield in OH groups on the composition of the precursor mixture.

Table 18.1 Absolute and relative yields of functional groups at the surface of pulsed plasma polymer layers measured with XPS after derivatization (cf. Section 18.2; 100 W).

Monomer

Theoretical stoichiometry

Measured stoichiometry

Yield in functional groups (per 100 C)

Characteristic side-products

Allyl alcohol Allylamine

100 C/33 O 100 C/33 N

100 C/33 O 100 C/36 N/4 O

31 (95%) 21 (65%)

Acrylic acid Ethylene

100 C/67 O 100 C

100 C/60 O 100 C/3 O

26 (79%) –

>C=O< C–N–C, C:N, C–O >C=O >C=O, C–O, C=C

18.3 Results

polyethylene-like layers were deposited in the same way. The nature and the quantity of the side-products of these polymers were investigated also. Generally, the number of side-products was minimal when low power was used (Table 18.1). 18.3.2 Surface Free Energy Measurements

The polar component increased exponentially with growing concentration of OH or COOH groups in the plasma copolymers of allyl alcohol–ethylene, allyl alcohol–butadiene or acrylic acid–butadiene precursors (Fig. 18.5). Low concentrations of OH or COOH groups produce only a very small polar component (Fig. 18.5). It is assumed that in this case the functional groups at the polymer surface are oriented toward the bulk, and thus do not contribute to the polar component. The measured surface energy of the pure ethylene homopolymer (plasma power input 300 W) was 36 mJ m–2, which is in the range of commercial polyethylene [1]. The polar component was also near zero, which qualifies the ethylene homopolymer as a pure chain-extending component in the copolymer and confirms the appropriateness of copolymers with ethylene sequences as a model surface with a variable concentration of exclusively one type of functional group. The dispersion component was determined to be 35 mJ m–2. Thus, the existing imperfections in the structure of the pulsed-plasma ethylene homopolymer (C=C double bonds, branched structures, and other inhomogeneities) did not influence the dispersion component noticeably. The pulsed-plasma polymerized poly(allyl alcohol) homopolymer possesses a surface energy of 51.6 mJ m–2. The butadiene homopolymer produced with 300 W has a surface energy of

Fig. 18.5 Polar component of surface energy of plasma polymers as a function of the number of functional groups in them.

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38.5 mJ m–2 and possesses a more highly polar component (8.5 mJ m–2) than the homopolymer formed at 100 W (3 mJ m–2), thus confirming the assumption that more regular and defined structures are produced under “softer” plasma conditions (here, lower power input). The dependence of polar components for the allyl alcohol–butadiene plasma copolymers (produced with 100 or 300 W) on the concentration of OH groups showed a greater slope in the exponential behavior for the 100 W copolymers than for the 300 W species, which is not yet well understood (cf. Fig. 18.5). 18.3.3 Peel Strength Measurements of Al-Plasma Modified PP Systems

Unspecifically oxygen-plasma modified PE and PP surfaces show maximum peel strength values respectively for a degree of oxidation of 20 O/100 C atoms (Fig. 18.6). No difference was measured between the Al peel strength of PP after specific or unspecific plasma pretreatment (Table 18.2). This was interpreted roughly by comparing the average of the polarities of different O functional groups pro-

Fig. 18.6 Al peel strength on PE, PP, and PET versus oxygen concentration introduced.

Table 18.2 Al peel strength of oxygen-plasma treated PP or oxygen-plasma treated + diborane-reduced PP.

Type of pretreatment

Measured peel strength [N m–1]

Scheme

O2 plasma

&450

O2 plasma + reduction

&450

CH COOH O CHO COH O /n j j jj j j CH CH CH CH CH CH j j j j j j

18.3 Results

duced by the O2 plasma (product spectrum: ca. 50% C–O–C, OH; 30% C=O, CHO; 20% COOH, COOR) with the polarity of OH groups. So it was assumed that in both cases similar polarities occur, as is confirmed by the similar contact angles of water. 18.3.4 Peel Strength of Al–Plasma-Produced Homopolymer–PP Systems

The pulsed plasma-initiated radical polymerization of monomers bearing functional groups was used to produce this kind of model surface. The peel strengths of Al on plasma homopolymer–PP systems are presented in Table 18.3. It was confirmed that layers of pulsed plasma polymerized ethylene homopolymer did not promote any adhesion to Al when applied in Al–PP systems as an adhesion-promoting interlayer. With NH2-containing layers, very weak peel forces were produced, in contrast to the high peel strength using OH- and especially COOH-containing interlayers (Table 18.3). In the case of CH2 and NH2 groups, failure occurs at the plasma polymer–PP interface, according to XPS of the peeled surfaces. With OH groups, the peel front propagates within the allyl alcohol plasma polymer. The adhesion of acrylic acid layers to PP as well as to Al was so strong that peeling occurs within the PP substrate, according to XPS analysis of both peeled surfaces. Thus, the interactions between aluminum and the single-type functional groups of homopolymers depend strongly on the type of functional group in the order: COOH OH NH2 > CH2 18.3.5 Peel Strength of Al–Plasma Copolymer–PP Systems

In the case of allyl alcohol–ethylene plasma copolymers, the maximum adhesion (650 N m–1) was measured at 27–29 OH per 100 C atoms. Therewith, this copolymer exhibits a significantly higher peel strength than the allyl alcohol homopolymer (85 N m–1). Thus, the above ranking of the adhesion-promoting

Table 18.3 Dependence of the Al peel strength on PP on the nature of the adhesion-promoting plasma homopolymer interlayer.

Functional group

Peel strength [N m–1]

Concentration

Type of failure

Location

CH2 NH2 OH COOH

0–10 &30 &85 &650

– 18 NH2/100 C 31 OH/100 C 24 COOH/100 C

interfacial interfacial cohesive cohesive

plasma polymer–PP plasma polymer–PP inside plasma polymer inside PP

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efficiency of different functional groups of homopolymers changes slightly if the peel strength of allyl alcohol–ethylene plasma copolymers is considered (cf. Fig. 18.7 and Table 18.4): COOH > OH NH2 > CH2 A nearly linear dependence of peel strength on the density of OH, NH2, and COOH functional groups was observed in the range 0–27 OH, 0–15 NH2, and 0–5 COOH per 100 C (see Fig. 18.7). Using poly(allyl alcohol), the peel front propagates along the plasma polymer–PP interface (interfacial failure). A plateau of maximum peel strength (650 N m–1) at concentrations of 27–29 OH groups/100 C atoms is observed. At these concentrations, partially or complete cohesive failures (inside the copolymer) occur at peeling. The pure allyl alcohol plasma polymer is tacky and weak and shows a low cohesive strength, resulting in low peel strength and a pure cohesive failure within the allyl alcohol homopolymer layer. NH2 groups also show a pure interfacial failure at the Al–plasma polymer interface and the lowest maximum peel strength at Al (see Table 18.4). COOH groups produce the highest peel strength at Al (see Fig. 18.7 and Table 18.4). In both cases of deposition of acrylic acid–ethylene or –butadiene

Fig. 18.7 Al peel strength of Al– plasma copolymer–PP composites as a function of functional group density at the plasma polymer surface (squares: COOH; half-filled squares: COOH; circles: OH; triangles: NH2; copolymers: ethylene and butadiene).

Table 18.4 Dependence of maximum Al peel strength on PP on the nature of the adhesion-promoting plasma copolymer interlayer.

Functional group

Concentration

Strength [N m–1]

Type of failure

Location

NH2 OH COOH

12–18 NH2/100 C 27–29 OH/100 C 8–24 COOH/100 C

10–30 600–620 &650

interfacial interfacial cohesive

plasma polymer–PP plasma polymer–PP inside PP

18.3 Results

pulsed-plasma copolymers (100 W, duty cycle 0.1), the Al peel strength depends linearly on the concentration of functional groups over the range 0–5 COOH/ 100 C. Higher concentrations of COOH groups (>5 COOH/100 C) do not increase the peel strength further. This plateau in peel strength at higher concentrations of COOH groups is characterized by peel specimens with pure cohesive failures. The interpretation is that the interactions between Al and COOH groups become too strong, and therefore the peel failure changes from a pure interfacial to a cohesive failure within the polypropylene substrate [22]. Also, this result corresponds well with the value of 15 O per 100 C atoms for similar treatments and systems given by Wu [1]. These values of 15–20 O-functional groups also correspond well with the concentrations of functional groups at the surface of plasma polymer interlayers needed to establish maximum adhesion. As discussed above, in this case 5–24 COOH or 27–29 OH groups/ 100 C atoms were needed to produce maximum peel strength (each 600– 650 N m–1). It must be remembered that in the case of O2 plasma treatment OH, COOH, and other O species are produced. The difference between these two types of adhesion-promoting plasma modifications results in the lower maximum peel strength in the case of O2 plasma treatment (400–450 N m–1) compared with the case of the deposition of plasma polymer interlayers (600– 650 N m–1). Moreover, with the combined process (O2 plasma and diborane treatment) not more than 14 OH/100 C could be produced – many fewer than the 30 OH/100 C of the allyl alcohol copolymer. Thus, the maximum peel strength should also be lower. Here, it can be assumed that the O2 plasma modification causes polymer degradation at the surface, thus forming a weak boundary layer. This was concluded from the abrupt decrease in peel strength when the O concentration of 20 O/100 C was exceeded. That points to a concurrent process, polymer degradation, which starts at the beginning of plasma exposure [74] and is caused by the vacuum UV radiation (k < 200 nm) and to a much lower extent by the plasma particle bombardment [75]. So, the adhesion-improving effects of the functional groups and the polymer degradation are superposed and prevent the maximum peel strength, as measured with the plasma, being reached. Moreover, in contrast to the pulsed plasma mode for polymer deposition, the O2-plasma functionalization was performed in the cw mode, thus increasing the irradiation dose of the polymer substrate. 18.3.6 Plasma Pretreatment of PTFE Surfaces

A few variants for producing well-adhering Al–PTFE systems were tested: (a) deposition of plasma polymers bearing a single type of functional group as adhesion promoters on virgin PTFE; (b) H2 plasma pretreatment of PTFE followed by deposition of an adhesion-promoting plasma polymer layer; and (c) H2 plasma pretreatment of PTFE alone. Process (a) did not give any Al–PTFE bond strength because of the poor adhesion of the plasma polymers to the PTFE surface. Process (b) provided adequate

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and strongly adhering Al–PTFE systems, whereas process (c) produced moderately to highly adhering Al–PTFE systems. Similarly to the results described in Refs. [76–78], defluorination in a continuous-wave or pulsed (0.1 ms “plasma-on”, 0.9 ms “plasma-off”) hydrogen RF plasma leads to a minimum of 17 F/100 C (35 F/100 C in Ref. [79]) and to 55 O/ 100 C (6.4 O/100 C in Ref. [79]) (Fig. 18.8). The introduction of oxygen is most likely due to post-plasma reactions of plasma-produced C-radical sites at the PTFE surface by reaction with molecular oxygen when the samples were transferred from the plasma reactor to the XPS spectrometer with transient contact with air [73, 80]. This post-plasma reaction is unavoidable, but in this case it may help to promote the adhesion between the deposited plasma polymer layer and the pretreated PTFE. It should be noted that the F/100 C ratio increases slightly and the O/C ratio (exsitu XPS measurements) increases significantly at treatment times longer than 120 s (Fig. 18.8). An explanation for this second increase in the F/C ratio may be the increased temperature at longer treatment times and, therefore, enhanced hydrophobic recovery [81]. On the other hand, a simple abrasion test (wiping with a cotton cloth) and the solvent stability test (6 h of rinsing in tetrahydrofuran– THF) of the defluorinated PTFE surface show that this near-surface layer became unstable when exposed for longer than 10 s to hydrogen plasma. The respective XP spectra became very similar to the spectrum of untreated PTFE. Therefore, for further plasma-chemical processing and metal deposition a H2-plasma treatment time of 10 s was chosen. In the bimodal shape of the C1s signals (CF2 and CHx), characteristic changes can be observed by prolonging the time of exposure to the hydrogen plasma. The broadening of that C1s peak at higher binding energies suggests the existence of C=C double bonds and O-bonded C species. This interpretation was also given in Refs. [1] and [82–85]. Hydrogen-plasma reduction of PTFE produces preferentially a CHx component with a C1s electron binding energy BE = 285 eV. Graphite-like species at BE = 284.5 eV could not be detected. There-

Fig. 18.8 Fluorine removal and oxygen incorporation as a function of treatment time in the H2 plasma (100 W, 6 Pa, cw mode).

18.3 Results

fore, it was concluded that a complete reduction of the PTFE backbone to loosely bonded graphite did not take place as sometimes occurs in the case of alkaline metal reduction. It can be speculated that hydrogen reduction at short treatment times ( 1000 s) results in slightly lower peel strength. Then, failure is observed between Al and PTFE.

18.3.7.2 Hydrogen Plasma Pretreatment of PTFE and Deposition of Plasma Polymer Layers The introduction of an adhesion-promoting pulsed plasma polymer interlayer onto the H2-plasma pretreated (each for 10 s) PTFE substrate improved the peel strength further to a range of 350–400 N m–1 limited by the adhesion of the adhesive (supporting) tape to the evaporated Al layer. Because of this limitation of the peel test, the measured peel strength is no measure for the difference between OH and COOH groups at the PTFE interface in this case (Fig. 18.9).

Fig. 18.9 Al peel strength as a function of pretreatment time and type of modification. Full symbols: interfacial failure between Al and PTFE; empty symbols: interfacial failure between tape and Al; half-filled symbols: mixed failure.

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However, the single-type functionalized PTFE interface gave a higher peel strength than the pure hydrogen-plasma modified PTFE without a plasma polymer interlayer. The locus of failure during peeling at maximum peel strength was identified by analyzing with XPS the peeled Al and PTFE surfaces of H2-plasma modified and pulsed plasma poly(allyl alcohol)-coated PTFE evaporated with Al (200 nm). At maximum peel strength, the peeled PTFE surface showed an XPS C1s signal at a binding energy of 291 eV characteristic of unmodified PTFE, and the peeled metal surface exhibited a signal near 285 eV that was assigned to CHx components. Thus, the peel front should have propagated along the interface between unmodified PTFE and the H2-plasma modified PTFE near-surface layer.

18.4 Discussion 18.4.1 Contribution of Chemical Bonds to the Resulting Adhesion Strength

The metal–polymer interactions, and therefore the peel strength, are based on the sum of the discrete chemical bonds between the aluminum and the functional groups, and other unspecific interactions as described above. For the interpretation of the observed effects of single-type functional groups on the peel strength of electron-beam evaporated Al layers, only the common considerations could be made in this work because of the lack of direct spectroscopic evidence for the interface chemistry. However, the strong differences between the low efficiency of CHx and NH2 groups on the one hand, and the high Al adhesionpromoting efficiency of OH and COOH groups on the other, correspond exactly to the chemical reactivity of evaporated Al toward these functional groups [28]. Hence, these chemical aspects should be explained in more detail. CH2 (ethylene-like) groups did not show any interaction with electron-beam evaporated aluminum – the peel strength is zero. Therefore, it was concluded that any chemical and physical interactions were absent. In the literature, contradictory results were obtained by considering the interactions between freshly evaporated Al atoms and aliphatic chains in polyethylene and polypropylene in more detail. The proposed formation of Al–C bonds did not influence the adhesion strength [30, 31] and could not be confirmed in other experiments [85, 86]. These Al–C bonds are unstable toward hydrolysis, and thus they should not contribute significantly to the peel strength [24]. One can argue that aliphatic plasma polymers from ethylene or propylene produced in the continuous-wave mode have significant structural anomalies in comparison with the respective classical products [24] and therefore they should increase the adhesion strength rather than keeping it at a constant (low) level. As seen in this work, the structural anomalies of aliphatic layers were minimized by the use of the pulsed plasma technique. This is demonstrated by the

18.4 Discussion

low surface energy as well as by the XPS data and the zero peel strength. Thus, it can be concluded that pulsed plasma polyethylene also does not interact with Al. Using primary amino groups, very low peel strength also was measured in this work. This is not surprising when one considers the possibilities of chemical interactions from the point of view of metal-organic chemistry. For the tentative formation of Al–N bonds the same limitations are valid as for Al–C bonds. Primary amino groups form Me–N bonds in contact with electronegative metals [33, 34]. When Al was evaporated onto model alkanethiol SAMs (self-assembled monolayers) with NH2 end-groups, the formation of Al–N bonds was observed [85]. However, the authors stated that impurities in the SAM layer (oxidation products, carbamate formation) contributed significantly to the Al–SAM interaction and always superimposed on the original Al–NH2 interactions. The Al– NH2 interactions may be of the Lewis type [87]. As shown here, the measured low peel strength of Al to NH2-modified polymer surfaces indicates weak interactions in the Al–NH2 system. In contrast, the strong reaction of metallic aluminum with OH groups is well known [33]. OH groups form Al–O–C bonds known as alcoholates [37]. The measured high adhesion strengths (see Tables 18.3 and 18.4) suggest that chemical interactions or covalent bonds play an important role. At this point, it should be remembered that CH2 groups alone did not improve the peel strength. Therefore, the adhesion improvement of OH-containing copolymers with aliphatic sequences should be caused exclusively by OH groups. In the same way, NH2-bearing plasma polymers were also not adhesion-improving (allylamine). However, when the NH2 groups were replaced by OH groups (allyl alcohol), the structure of both polymers should be comparable and then the peel strength increases. Thus, the conclusion seems to be possible that the Al–OH interactions exclusively determine the adhesive strength. From the viewpoint of metal-organic chemistry the most important Al–OH interactions are in alcoholates, as mentioned above. This was evidenced by the extensive consumption of evaporated Al in interfacial bond formation [40]. Metallic Al was only detected when thick Al layers were evaporated [28]. Thus, all the OH groups at the surface are consumed by reaction with Al atoms [85]. Generally speaking, Al evaporated onto polymer surfaces always reacts with O present in C=O, C–O–C, and OH groups. However, it can be speculated that the Al–O–C bonds are only an intermediate state. From the physical chemistry point of view, it seems to be plausible that the high electronegative redox potential of Al (–1.51 V) and, therefore, its affinity to oxygen is so high that the endpoint of interfacial reactions is the redox reaction of OH groups and Al with formation Al(OH)3 and Al2O3 [40, 46–48]. The consequence of this redox reaction for the adhesive strength is not known. COOH groups interact strongly with metallic Al with the result that all the carboxylic groups are consumed [85]. It is the same situation as with OH groups. First, Al reacts with the acidic groups with the formation of salt-like structures [Eq. (C)].

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2 Al 2 R COOH ! 2 Al COO H2

C

Moreover, Al can also form weak interactions with COOH groups: this happens preferentially to the C=O site first, thus forming A C–O interactions [88] but there follows the formation of (weak) interactions to the C–O site [46]. The end product of the Al–COOH reaction is Al2O3. Its formation can be considered as a redox reaction again [40, 84] because of the reduction of C=O and C–O sites and the oxidation of Al. In this case, unspecific thermodynamic adhesion must be implied. The exact experimental proof of all the interactions and bond formation discussed here is matter for future work.

18.4.2 Dependence of Adhesion Strength on Concentration of Functional Groups at the Polymer S

Interpreting the peel strength dependence on the concentration of functional groups, the different slopes of the linear section of the graph, starting from the origin, reflect the strength of the respective metal–polymer bonds (see Fig. 18.7). Thus, about 5 COOH and 27 OH per 100 C are needed to reach > 600 N m–1 peel strength (the maximum peel strength before cohesive failure appears), whereas NH2 or CH2 groups allow, at best, 30 N m–1. These results correspond very well with those obtained with metal–polymer systems where functional groups are present in the polymer matrix [1]. For example, systems of Al with poly(vinyl chloride)–maleic acid copolymers were produced [89, 90]. The authors synthesized copolymers with variable contents of COOH groups. The dependence of the lap shear strength of Al on the COOH content in the chemically synthesized copolymers shows the same type of peel curve as in the case of the Al–plasma polymer–PP systems. Also in this case the joint strength increases rapidly with the concentration of COOH and then levels off. The authors found a curve progression which is described by Eq. (3), where rf,m is the fracture strength at complete surface coverage, h the fractional surface coverage, c the concentration of functional groups, and b is a constant [1]. rf rf ;m h rf ;m bc=1 bc

3

The authors argue that the plateau in the peel strength may arise either from the saturation adsorption of the functional groups at the interface, or from cohesive failure of the adhesive or adherend [89, 90]. Formally, there are similarities in curve progressions between this chemical bulk modification with COOH groups [1, 89, 90] and our work, with surface modification by COOH-containing adhesion-promoting layers. In this work, the adhesion strength levels off at higher COOH concentrations because of the adhesion-limiting cohesive failure in the PP substrate.

18.5 Conclusion

Considering the peel strength results presented in Figs. 18.7 and 18.9, the following relative ranking can be given for the efficiency of the respective Al–X interactions in the Al–PP system: COOH: OH: NH2 : CH2 130 : 24 : 1 : 0 One problem remains regarding the deposition of Al layers onto model surfaces with very low concentrations of (single-type) functional groups, and must be investigated in the future. This problem is concerned with the mode of metal film formation, i.e., whether homogeneous film formation or island growth is obtained (see Ref. [91]). This behavior may have an influence on the resulting peel strength. However, the linearity of the peel strength versus concentration of functional groups (in the ranges 0–5 COOH, 0–27 OH, and 0–18 NH2 per 100 C) gives a hint that no significant influence of film formation processes exists. In the case of the mechanism of film formation changing at low concentrations of functional groups, one would expect the appearance of a point of discontinuity.

18.5 Conclusion

Adhesion-promoting plasma polymer interlayers, possessing a maximum of 31 OH or 21 NH2 or 25 COOH groups per 100 C atoms, were produced by applying low power and using the pulsed plasma technique. By carrying out a pulsed plasma-initiated but chemically dominated copolymerization, the density of single-type functional groups could be varied continuously between 0 and 31 OH, 0 and 21 NH2, and 0 and 25 COOH groups per 100 C atoms. With such model layers bearing functional groups as adhesion-promoting interlayers in Al–PP systems, the peel strength of evaporated aluminum layers increases in the order: CH2 < NH2 OH < COOH This ranking of adhesion promotion should be attributed to the ability of the functional groups to form chemical bonds with aluminum. With carboxylicmodified surfaces, the adhesion at the Al–COOH interface is greater than the cohesive strength of the polypropylene substrate. The peel strength depends linearly (see Fig. 18.7, 0–5 COOH, 0–27 OH, and 0–18 NH2 per 100 C) on the density of functional groups. Similar results were obtained when Al was evaporated onto PTFE as substrate. However, in this case, the PTFE must first be defluorinated in the near-surface layer using H2 plasma before it is coated with pulsed plasma polymer layers as an adhesion promoter for the evaporated Al. For PTFE the ranking is: H2 plasma < OH < COOH

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19 Alkene Pulsed Plasma Functionalized Surfaces: An Interfacial Diels-Alder Reaction Study F. Siffer, J. Schultz, and V. Roucoules

Abstract

The kinetics of an interfacial Diels-Alder reaction on maleic anhydride pulsed plasma polymer thin films has been investigated. This entails the aminolysis reaction of allylamine with maleic anhydride pulsed plasma polymer films to yield terminal alkene groups (dienophile groups) at the surface. Subsequent exposure to [(trimethylsilyl)methyl]cyclopentadiene leads to a Diels-Alder cycloaddition reaction. The kinetics of the reaction was followed by contact angle measurements with the application of the Cassie equation. Thermodynamic parameters, such as the activation energy determined according to the Arrhenius equation and the activation entropy, were calculated according to the transition-state theory and discussed. The reactivity of dienophile groups confined in pulsed plasma polymer thin films is compared with the behavior of dienophile groups in monolayers.

19.1 Introduction

Pulsed plasma polymerization has become a useful candidate in all adhesion problems dealing with functionalized surfaces [1]. Some inherent advantages of this approach include the fact that the plasmachemical surface functionalization step is substrate-independent [2]. Pulsed plasma polymerization entails modulating the electrical discharge on the millisecond–microsecond timescales. In the case of gaseous precursors containing polymerizable carbon–carbon double bonds, pulsed plasma functionalization comprises two distinct reaction regimes corresponding to the plasma duty cycle on- and off-periods. This gives rise to extremely high levels of structural retention and incorporation of specific functional groups at the surface [3–12]. By programming the pulse duty cycle, it is possible to tailor the surface density of the desired functional groups. Recently, this technique has been successfully applied to obtain alkene-functionalized surfaces [13]. These surfaces have been evaluated for their suitability Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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as dienophiles for the Diels-Alder cycloaddition reaction with 1,3-cyclohexadiene. However, no study has been done to elucidate the dependence of the rate of the interfacial Diels-Alder reaction on the steric environment of the plasma-immobilized dienophiles. The understanding of this dependence becomes essential when the aforementioned surfaces are involved in adhesion problems. In this study, the behavior of dienophile groups in monolayers is compared with the reactivity of these dienophile groups confined in pulsed plasma polymers.


Maleic anhydride (Prolabo, 99.5% purity) was ground into a fine powder and loaded into a stoppered glass gas-delivery tube. Plasma polymerization experiments were carried out in an electrodeless cylindrical glass reactor (6 cm diameter, 680 cm3 volume, base pressure 5 ´ 10–4 mbar, and with a leak rate better than 1.0 ´ 10–10 kg s–1) enclosed in a Faraday cage. The chamber was fitted with a gas inlet, a Pirani pressure gauge, a two-stage rotary pump (Edwards) connected to a liquid nitrogen cold trap, and an externally wound copper coil (4 mm diameter, five turns). All joints were grease-free. An matching network (Dressler, VM 1500 W-ICP) was used to match the output impedance of a 13.56 MHz RF power supply (Dressler, Cesar 133) to the partially ionized gas load by minimizing the standing wave ratio of the transmitted power. During electrical pulsing, the pulse shape was monitored with an oscilloscope and the average power hPi delivered to the system was calculated from the expression: hPi = Pp[ton/(ton + toff)], where Pp is the average continuous-wave power output and ton/(ton + toff) is defined as the duty cycle. Prior to each experiment, the reactor was cleaned by scrubbing with detergent, rinsing in propan-2-ol, and ovendrying; this was followed by a 30 min high-power (60 W) air plasma treatment. The system was then vented to air and a gold-coated glass slide was placed in the center of the chamber before evacuation back down to base pressure. Subsequently, maleic anhydride vapor was introduced into the reaction chamber at a constant pressure of 0.2 mbar and with a flow rate of approximately 1.6 ´ 10– 6 g s–1. At this stage, the plasma was ignited and run for 30 min. The optimum deposition conditions correspond to: power output = 5 W, pulse on-time = 25 ls, off-time = 1200 ls. These parameters have previously been optimized on the basis of a full factorial design and a central composite approach [14]. Upon completion of deposition, the RF generator was switched off, and the monomer feed allowed to continue to flow through the system for a further 2 min before venting up to atmospheric pressure. Next, reaction of the maleic anhydride plasma polymer deposited films with allylamine (Aldrich, 99+%) was carried out under vacuum without exposure to air. At this stage, timing of the surface functionalization reaction commenced. Upon termination of exposure, the allylamine reservoir was isolated, and the


whole apparatus pumped back down to the system base pressure. Then the alkene-functionalized surface was removed from the reactor and placed in an oven at 120 8C for 2 h in order to transform the amide groups into more stabilized cyclic imide groups. Next, the sample was immersed in a flask containing cyclopentadiene for 1 h or [(trimethylsilyl)methyl]cyclopentadiene for kinetics experiments. Afterward, the functionalized polymer surface was washed with dichloromethane and dried under a stream of nitrogen prior to the characterization. Alkene-functionalized monolayers were prepared as follows. Gold substrates were obtained by successive evaporation under high vacuum of 50 nm chromium and 100 nm gold onto clean glass slides. 16-Mercaptohexadecanoic acid was then adsorbed overnight at room temperature from a freshly prepared 2 mm ethanolic solution [15]. Cleaned substrates of the carboxylic acid SAMs (self-assembled monolayers) were then immersed for 20 min at room temperature in a freshly prepared solution of 0.1 m trifluoroacetic anhydride and 0.2 m triethylamine in anhydrous N,N-dimethylformamide in order to obtain anhydride groups [15]. The substrates were removed from the solution, rinsed thoroughly with dichloromethane, and dried in a nitrogen stream before analysis. The next steps were similar to those for the plasma polymer, i.e., reaction in the gas phase with allylamine followed by a 2 h annealing step at 120 8C to obtain cyclic imide groups. Cyclopentadiene was obtained by transferring dicyclopentadiene into a 250 mL flask and heating until boiling (the temperature of the vapor must not exceed 50 8C). In order to obtain [(trimethylsilyl)methyl]cyclopentadiene, a mixture of (chloromethyl)trimethylsilane (Aldrich) and sodium cyclopentadienide in anhydrous tetrahydrofuran (Aldrich) was stirred for two days at room temperature under a nitrogen atmosphere. The dark red liquid obtained was filtered and distilled at room temperature (0.15 torr). The final product was employed during further kinetic studies. Polarization-modulation infrared reflection–absorption spectroscopy (PM-IRRAS) spectra were recorded with a Bruker IFS 66/S Fourier transform infrared spectrometer equipped with a PMA 37 polarization modulation module and a nitrogen-cooled MCT detector. The infrared beam was first p-polarized with a ZnSe wire grid polarizer (Specac) before passing through a photoelastic modulator (Hinds Instruments, PEM-90), which modulated at a frequency of 74 kHz. A lock-in amplifier (Stanford model SR-830) was used to obtain the PM-IRRAS spectra. The half-wave retardation frequency was set at 4000 cm–1. The PM-IRRAS spectra were recorded as S = (Rp–Rs)/(Rp + Rs). A total of 250 scans at a resolution of 4 cm–1 were collected for each measurement at an angle of incidence of 82.58 with respect to the normal to the sample surface. Sessile drop contact angle measurements were carried out using a contact angle measuring system G2 from Krüss with 3 lL high-purity water drops. The droplet was then enlarged or shrunk while measuring the contact angle to obtain the advancing contact angle. The volume of the drop was incremented dif-

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ferentially until the contact line was observed to advance. The contact angle obtained just before the meniscus moved was measured as the advancing contact angle. Surface roughness measurements were made with a Dimension 3000 scanning probe microscope (Digital Instruments). A silicon cantilever with a silicon tip was used for all measurements. The spring constant of the cantilever was 20–100 N m–1 and the average tip radius was 20 nm. Typically, the surface morphology of 2 lm ´ 2 lm spots near the center of each sample was observed in the tapping mode of the scanning probe microscope. Film thickness values were estimated by ellipsometry. The ellipsometry measurements were performed on a phase modulation Multiskop from Physik Instrumente (Model M-033k001) at 632.8 nm (He–Ne laser). The cross-section of the laser beam was about 1 mm2. Measurements in air were performed at different positions (at least five) of the samples in order to check the quality of the films.

19.3 Results and Discussion 19.3.1 Interfacial Chemistry

Smooth thin films of maleic anhydride polymer with a thickness of 15.9 ± 0.8 nm were obtained by plasma polymerization. The tapping-mode atomic force microscope (AFM) image acquired in air showed almost featureless films with a root-mean-square (rms) roughness of 0.19 ± 0.02 nm assessed from 4 lm2 images (Fig. 19.1). The height image shows a relatively smooth surface with particles of approximately 30 nm diameter, whereas the phase image shows small chemical heterogeneities with a contrast of 28. This low roughness is a typical result of the competition between the polymerization mechanisms which occur at the same time in the gas phase and on the surface [16, 17]. Infrared analysis of the deposited maleic anhydride pulsed plasma polymer films confirmed a high degree of anhydride group incorporation. The following characteristic infrared absorption features of cyclic anhydride groups were identified: asymmetric and symmetric C=O stretching (1860 and 1796 cm–1), cyclic conjugated anhydride group stretching (1241–1196 cm–1), C–O–C stretching vibrations (1097–1062 cm–1), and cyclic unconjugated anhydride group stretching (964–906 cm–1) (Fig. 19.2, spectrum a). The reaction of allylamine (Fig. 19.2, spectrum b) from the gaseous phase with the deposited maleic anhydride pulsed plasma polymer layer resulted in a ring opening of the cyclic anhydride centers to yield the corresponding amide (amide I at 1658 cm–1 and amide II at 1550–1510 cm–1), carboxylic acid stretching (1716 cm–1) infrared bands (Fig. 19.2, spectrum c). Heating of these surfaces to 120 8C for 2 h gave rise to the formation of cyclic imides as seen by the drop in intensity of strong amide


Fig. 19.1 (a) Tapping mode AFM topographic image of the optimized maleic anhydride pulsed plasma polymer thin film on a silicon wafer; (b) the corresponding phase image.

and C=O acid bands, while two imide stretching vibration bands appear at 1775 and 1710 cm–1 (Fig. 19.2, spectrum d). Other new features include CN stretching vibrations at 1337 cm–1 and two new bands at 1210–1140 cm–1 characteristic of CNC symmetric and asymmetric stretching (imide stretching) respectively. However, no changes were observed in the C=C stretching (1638 cm–1) and the H–C = stretching (995 cm–1) bands associated with the terminal alkene groups. Other new spectral features include very strong bands centered at 2930 cm–1 attributable to the characteristic asymmetric CH2 stretching and at 2850 cm–1 attributable to the symmetric CH2 stretching (Fig. 19.3, spectrum a). The weak band centered at 3030 cm–1 is associated with asymmetric = CH2 stretching while the weak band at 2995 cm–1 correlates with the symmetric = CH2 stretching mode of the alkene groups (Fig. 19.3, spectrum a). All the infrared spectroscopic results were confirmed by X-ray photoelectron spectroscopy (XPS) analysis (results not shown).

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19 Alkene Pulsed Plasma Functionalized Surfaces: An Interfacial Diels-Alder Reaction Study Fig. 19.2 Infrared spectra of (a) the maleic anhydride pulsed plasma polymer thin film; (b) the allylamine in the gas phase; (c) the plasma polymer after reaction with allylamine; (d) the same surface as for (c), annealed at 120 8C for 2 h. CH2 scis = scission.

19.3.2 Cycloaddition

Other newly appearing absorbances include strong bands centered at 2956 cm–1 attributable to the characteristic cyclic asymmetric CH2 stretching and at 2875 cm–1 attributable to the cyclic symmetric CH2 stretching (Fig. 19.3, spectrum b). A control experiment entailed immersion of the cyclic imide functionalized plasma polymer surface into THF at 25 8C for 1 h. No changes in the infrared spectrum were observed (not shown). The intermediate amide functionalized plasma polymer surface was also exposed to a solution of cyclopentadiene in THF at 25 8C for 1 h. Infrared analysis showed spectral features similar to those described above for the imide surface, the main difference being the peak between 1800 and 1600 cm–1 indicating the presence of amide rather than imide linkages at the surface. Finally, the plasma polymer surface functionalized with cyclic imide groups was exposed to [(trimethylsilyl)methyl]cyclopentadiene solution in cyclohexane at 25 8C for 1 h. Two new bands appeared at 2975 and 2890 cm–1 characteristic of the asymmetric CH3 stretching and the symmetric CH3 stretching (Fig. 19.3, spectrum c).

19.3 Results and Discussion Fig. 19.3 Infrared spectrum in the 3200 cm–1–2800 cm–1 region of (a) the cyclic imide functionalized plasma polymer; (b) the same polymer as for (a), after Diels-Alder reaction with cyclopentadiene (c) the same polymer as for (a) after Diels-Alder reaction with [(trimethylsilyl)methyl]cyclopentadiene.

19.3.3 Kinetics

A comparative study of the temperature dependence of the Diels-Alder reaction between [(trimethylsilyl)methyl]cyclopentadiene and dienophile groups confined in self-assembled monolayers or in pulsed plasma polymer layers has been done. The reactivity of dienophile groups confined in pulsed plasma polymer thin films is compared with the behavior of dienophile groups in monolayers, because of their well-known arrangement properties.

19.3.3.1 Monolayers The structure of alkene-functionalized monolayers was characterized by contact angle measurements. An advancing contact angle of 74 ± 28 was observed. The progress of the Diels-Alder reaction between [(trimethylsilyl)methyl]cyclopentadiene (TMSM-Cp) and dienophile groups confined in monolayers can be followed by contact angle measurements performed at 25 8C for various Diels-Alder reaction temperatures, as shown in Fig. 19.4. For increasing reaction temperatures, the advancing contact angle increases more rapidly, indicating a faster rate of reaction. The surface coverage with bicyclic groups was calculated

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Fig. 19.4 Advancing water contact angles, measured at 25 8C for monolayers after Diels-Alder reaction at different temperatures with TMSM-Cp, versus reaction time.

from the data using the Cassie equation [Eq. (1)], where vC=C and vbicyclic are the surface coverages of the two components and hC=C and hbicyclic are the advancing contact angles of the two pure monolayers. cos hexp vCC cos hCC vbicyclic cos hbicyclic

1

In Fig. 19.5 the inverse of the corresponding coverage is plotted versus the time of reaction. For all temperatures studied, we observed a linear dependence. This linear behavior indicates that the Diels-Alder reaction proceeds homogeneously and can be described by standard pseudo-second-order kinetics. The rate of the reaction V is described by Eq. (2), where [CpdSi] and vC=C denote the concentration of [(trimethylsilyl)methyl]cyclopentadiene and the density of alkene groups at the surface respectively. p and q represent the partial orders of the reaction and kDA the Diels-Alder rate constant. q

V kDA CpdSip vCC

2

Because (a) p = 1 (this result has been demonstrated in other papers [18] but has also been verified in our Laboratory) and (b) the experiments were carried out under pseudo-second-order conditions (the amount of diene is in large excess over the immobilized alkene groups and hence [CpdSi] & constant), the apparent rate constant (or the pseudo-second-order rate constant) k'' is equal to kDA[CpdSi]. The slopes of the least-squares fits of the linearized data shown in Fig. 19.5 correspond to the pseudo-second-order rate constant k'' of the Diels-Alder reaction at the respective temperature. These concentration-dependent rate con-


Fig. 19.5 Linearization according to pseudo-second-order kinetics of the alkene group surface coverage on monolayers. The solid lines correspond to least-squares fits of the data.

stants can be converted to concentration-independent Diels-Alder rate constant kDA (Table 19.1). The Diels-Alder rate constants obey the Arrhenius equation [Eq. (3)], as shown in Fig. 19.6. From the slope of the linearized plot, an activation energy Ea of 71.6 ± 2.8 kJ mol–1 can be estimated. ln kDA ln A

Ea =RT

3

By applying transition-state theory, we can calculate the activation entropy DS= of this Diels-Alder reaction from Eq. (4), where R is the molar gas constant, A is the preexponential factor in the Arrhenius equation, T is the absolute temperature, kb is the Boltzmann constant, and h is Planck’s constant. The values of Ea and DS= are presented in Table 19.2. DS6 R lnA=T

ln kb =h

Table 19.1 Diels–Alder rate constant kDA calculated for the different substrates.

Sample

kDA [(m3 mol–1)2 s–1] a)

Monolayer Plasma polymer thin film

4.7904 ´ 10–10 3.5928 ´ 10–10

a)

Calculated for 295 K.

4

297

298


Fig. 19.6 Diels-Alder rate constants obtained from Fig. 19.5, plotted according to the Arrhenius equation. The solid line corresponds to a least-squares fit of the data.

Table 19.2 Activation energies and estimated parameters characterizing the transition state.

Sample

Ea [kJ mol–1]

A [(m3 mol)2 s–1]

DS a) [kJ mol–1]

DH a) [kJ mol–1]

DG a) [kJ mol–1]

Monolayer

71.7 ± 2.8

1967.3 ± 1549.6

–457.5 ± 12.9

69.2 ± 2.8

205.5 ± 1.1

Plasma polymer thin film

48.1 ± 3.7

0.12 ± 0.1

–538.2 ± 16.4

45.6 ± 3.7

206.0 ± 1.2

a)

Calculated for T = 298 K.

19.3.3.2 Plasma Polymer Thin Films The progress of the Diels-Alder reaction was assessed by contact angle measurements performed at room temperature (Fig. 19.7). Again, the reaction was studied systematically at six different temperatures. We observed that the Diels-Alder reaction could be described as a pseudo-second-order reaction (Fig. 19.8). Similarly to the Diels-Alder reaction on monolayers, the third-order rate constants kDA, calculated from the least-squares fits shown in Fig. 19.8 for the Diels-Alder reaction in the polymer thin film, obey the Arrhenius equation (Fig. 19.9). The activation energy Ea = 48.1 ± 3.7 kJ mol–1 and the activation entropy DS= = –538.2 ± 16.4 J mol–1 K–1 at 298 K (Table 19.2) are determined at the polymer surface in the same way as for the monolayers.


Fig. 19.7 Advancing water contact angles, measured at 25 8C after Diels-Alder reaction at different temperatures with TMSM-Cp for pulsed plasma polymer thin films, versus reaction time.

Fig. 19.8 Linearization according to pseudo-second-order kinetics of the alkene group surface coverage on pulsed plasma polymer thin films. The solid lines correspond to least-squares fits of the data.

19.3.3.3 Comparison of Surface Reaction in Monolayers and Plasma Polymer Thin Films Both monolayers and plasma polymer thin films studied show pseudo-secondorder surface Diels-Alder kinetics and obey the Arrhenius equation. The magnitudes of the rate constants kDA calculated in the case of the Diels-Alder reaction on pulsed plasma polymer thin films are lower than the magnitudes of the rate constants on monolayers (Table 19.1). The rate constants seem to reflect the

299

300


Fig. 19.9 Diels-Alder rate constants obtained from Fig. 19.8, linearized according to the Arrhenius equation. The solid line corresponds to a least-squares fit of the data.

Scheme 19.1 The confinement effect. The arrows illustrate the diffusion of the reactant (a) at the surface of the monolayer and (b) into the alkene-functionalized plasma polymer thin film.

“accessibility” of the alkene groups involved in the reaction: The reactivity of the alkene groups seems to be lower at the surface of the pulsed plasma polymer thin films than that of the alkene groups in the monolayers (Scheme 19.1). But we must note that this observation may also be attributed to a lower interfacial effective concentration of [(trimethylsilyl)methyl]cyclopentadiene or a lower density of alkene groups in the region of the pulsed plasma polymer film (the plas-


ma polymer used in this work contains only 32% anhydride centers besides other polar groups such as alcohols, acids, esters etc.). The activation energies and the parameters characterizing the transition state of the activated “complex” are calculated according Eqs. (5) and (6) and are summarized in Table 19.2 for the reactions on monolayers and on plasma polymer thin films as well. DH Ea DG DH

RT T DS

5 6

The activation energy Ea is significantly lower for the pulsed plasma polymer thin films than for the monolayers. This is understandable, since the surface region of the pulsed plasma polymer thin films can be expected to be less mobile (they have less mobility and flexibility) than the monolayers: Monolayers are not tightly packed SAMs whereas the structure of the pulsed plasma polymer film is more or less crosslinked. In plasma polymer thin films fewer degrees of freedom are available than in the monolayers and the vibrations are hindered by crosslinking effects. This results in a reduced activation entropy, which is the dominating effect for free activation enthalpy in our case. An increase in temperature will only lead to a modest increase in the rate constant whereas the monolayers are probably more open structures for which an increase in temperature will lead to a significant increase in the rate constant. Another explanation may be the difference in chemical environment between alkene groups confined in monolayers and alkene groups confined in plasma polymers. The polar groups in the plasma polymer layer may facilitate the reaction by lowering the activation energy Ea. The experiments show that the enthalpies and entropies of activation were very different for the two cases and that these differences compensated one another to give similar free energies of activation. The interfacial Diels-Alder reaction with the plasma polymer thin films had an enthalpy of activation DH= = 45.6 kJ mol–1 whereas the interfacial Diels-Alder reaction with monolayers had a DH= of 69.2 kJ mol–1, a value that is 23.6 kJ mol–1 higher than that for the pulsed plasma polymer thin films. This reflects the results discussed below for the activation energy and could be explained in the same way. In both cases, the negative values of DS= reveal that the transition-state complexes for the products are more ordered than the reactants, which is due to a specific alignment requirement of the p-system during the Diels-Alder reaction. A high negative value of DS= is unfavorable for the rate of reaction. The smaller absolute value of DS= for reaction of the alkene in monolayers requires that the transition-state complex is less ordered. This point is consistent with the larger DH= that accompanies the reaction on monolayers. All these data are consistent with differences in structure between monolayers and pulsed plasma polymer thin films. These differences in structure appear to be intimately linked to the kinetic and thermodynamic parameters of the reactions and the corresponding transition state. But we must note that these explanations do not

301

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address roles that the solute molecules can play near the interface, which could easily dominate changes in DS= [18]. Furthermore, we do not forget that we have limited experience of accessing interfacial properties [18]. It is not yet possible to fully interpret activation parameters for interfacial reactions [18, 19]. In future studies, we propose to add to the contact angle measurements (which probe only 5–10 Å of the layer) XPS and FTIR spectroscopy analysis, in order to understand the kinetics of the reaction in the interior of the pulsed plasma polymer thin film. Once quantitative elucidation of the reactivity of the pulsed plasma polymer thin film has been fully accomplished, adhesion strength measurements will be performed and correlations between adhesion parameters and thermodynamic parameters will be explored. This will be the subject of a further paper.

19.4 Conclusion

Bicyclo[2.2.1]hept-2-ene groups can be introduced onto solid substrates via the Diels-Alder type of cycloaddition reaction between cyclopentadiene or [(trimethylsilyl)methyl]cyclopentadiene and a well-adhering maleic anhydride pulsed plasma polymer layer which has been functionalized with dienophile groups. This approach offers great flexibility for tailoring the density of such cyclic rings at the surface, since one of the distinct advantages of pulsed plasma chemistry is that the number of reactive sites can be easily controlled by simply programming the pulsed plasma duty cycle. The quantitative elucidation of the surface confinement effect of dienophile groups on the Diels-Alder reaction led to the conclusion that the reaction at the pulsed plasma polymer surface is significantly different from the reaction in the monolayer. The plasma polymer thin films are less reactive than the monolayers but the transition-state complex is more ordered. This means that this transition-state complex is more stable at the pulsed plasma polymer surface than on monolayers because of the chemical environment of the molecules. Since the reaction at the plasma polymer is significantly more confined than in monolayers, the reaction is less affected by the temperature. These first results need to be completed by XPS and FTIR spectroscopic analysis in order to obtain quantitative elucidation of the reactivity in the entire pulsed plasma polymer thin film. This work also shed light on interesting properties of thin films due to the thermally reversible chemistry of the Diels-Alder reaction. Carefully designed, alkene- and diene-functionalized, pulsed plasma polymer thin films will make it possible to control adhesion between any kinds of solid surfaces.

References

References 1 M. Strobel, C. S. Lyons, K. L. Mittal, Eds.

2 3 4

5

6

7 8 9

10

Plasma Surface Modification of Sample Preparation. Polymers: Relevance to Adhesion; VSP: Utrecht, The Netherlands, 1994. H. Yasuda, Plasma Polymerization; Academic Press: London, 1985. L. M. Han, K. Rajeshwar, R. B. Timmons, Langmuir 1997, 13, 5941–5950. N. M. Mackie, N. F. Dalleska, D. G. Castner, E. R. Fisher, Chem. Mater. 1997, 9, 349–362. J.-H. Wang, X. Chen, J.-J. Chen, J. G. Calderon, R. B. Timmons, Plasmas and Polymers 1997, 2, 245–260. M. T. van Os, B. Menges, R. Foerch, G. J. Vancso, W. Knoll, Chem. Mater. 1999, 11, 3252–3257. C. Tarducci, S. Brewer, C. Willis, J. P. S. Badyal, GB Patent Appl. No. 0001514.9. S. J. Hutton, J. P. S. Badyal, GB Patent Appl. No. 9726807.2. C. Tarducci, S. Brewer, C. Willis, J. P. S. Badyal, Chem. Mater. 2001, 13, 1800– 1803. S. R. Coulson, S. Brewer, C. Willis, J. P. S. Badyal, GB Patent Appl. No. 9712338.4.

11 M. E. Ryan, A. M. Hynes, J. P. S. Badyal,

Chem. Mater. 1996, 8, 37–42. 12 S. Schiller, J. Hu, A. T. A. Jenkins, R. B.

13

14

15

16 17

18 19

Timmons, F. S. Sanchez-Estrada, W. Knoll, R. Förch, Chem. Mater. 2002, 14, 235–242. V. Roucoules, C. A. Fail, W. C. E. Schofield, D. O. H. Teare, J. P. S. Badyal, Langmuir 2005, 21, 1412–1415. F. Siffer, A. Ponche, P. Fioux, J. Schultz, V. Roucoules, Analytica chimia acta 2005, 539, 289–299. L. Yan, C. Marzolin, A. Terfort, G. M. Whitesides, Langmuir 1997, 13, 6704– 6712. E. Amanatides, S. Stamou, D. Mataras, J. Appl. Phys. 2001, 90, 5786–5798. C. I. Butoi, M. L. Steen, J. R. D. Peers, E. R. Fisher, J. Phys. Chem. B 2001, 105, 5957–5967. Y. Kwon, M. Mrksich, J. Am. Chem. Soc. 2001, 124, 806–812. A. J. Pertsin, M. Grunze, Langmuir 2000, 16, 8829–8841.

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20 Laser Surface Treatment of Composite Materials to Enhance Adhesion Properties Q. Bénard, M. Fois, M. Grisel, and P. Laurens

20.1 Introduction 20.1.1 Why Treat a Composite Surface?

In many industries and more particularly for aeronautical applications, composite materials are often used for various purposes such as obtaining low-weight properties when compared with conventional materials. However, bonding of two composites by classical techniques such as bolting, riveting, or screwing can involve a dramatic addition to the weight of such assemblies. Alternatively, bonding composites with an adhesive layer makes it possible to obtain lighter structures. However, in this case the composite surface needs specific care in order to lead to a strong and durable bond. As a composite surface is often considered as a polymer entity, most of the available surface treatments performed on it are identical to those used for polymer surfaces. Any surface treatment mainly aims to improve a specific surface parameter such as roughness, chemical composition, etc., as depicted in the various adhesion theories. 20.1.2 Available Treatments for Composite Surfaces

This section summarizes the surface treatments available and performed nowadays, with the adhesion theory they tend to improve. Most of the surface treatments aim to bring roughness on the adherend, in order to satisfy the mechanical adhesion theory proposed by McBain [1], which is mainly aimed to enhance the “mechanical interlocking” of the adhesive with the adherend. To this end, solid abrasion remains one of the most widespread techniques; solid particles are directed onto the composite surface using different particle sizes at controlled impact speeds. This kind of treatment makes it possible Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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20 Laser Surface Treatment of Composite Materials to Enhance Adhesion Propertiess

to remove contaminated layers and the roughened surface provides mechanical interlocking with the adhesive. Moreover, it is sometimes argued that the increased roughness forms a larger effective surface area for the bond [2, 3]. In other cases, solvents can also be used to etch the surface. These solvents are selected from strong acids or bases, and also oxidant solutions [4]. In addition, other techniques that are more specific to the composite domain are available for surface roughening. The peel ply or tear ply treatment [3, 5] consists of an impregnated superficial ply removed prior to bonding, thus leading to increased surface roughness, to cleaning, and to the removal of mold release agents. The resulting surface morphology is directly linked to the pattern of the peel ply used. Another theoretical aspect of adhesion can be optimized with surface treatment. The interdiffusion theory described by Voyutski [6] involves the diffusion of macromolecular chains across a polymer interface. In this case the interface is transformed progressively to a wide interphase. Welding of composites is then the main technique that can enhance this particular aspect. This welding can be performed with ultrasonic, electric induction, or resistance techniques [7]. The last two theories described (chemical and thermodynamic) are intimately linked together because both of them induce a modification of the chemical composition at the surface. On the one hand, this modification can change the thermodynamic parameters (wettability) of the surface. On the other, changes in chemical composition influence the chemical adhesion established between the adherend and the adhesive layer. Numerous treatments are available for surface modification: with coronas [8], plasmas [9, 10], lasers [11, 12], ion-assisted reactions [13], or coupling agents [14, 15]. All these treatments do not only change the chemical composition; they can also affect the roughness, the orientation of macromolecular chains, and the mechanical behavior. In order to understand the aim of surface treatments for composite bonding, we will concentrate on a real composite bonding problem for aeronautical purposes. Since classical surface treatments like peel ply can be limited by a cohesive failure occurring in the material, we will focus on a new kind of surface treatment (excimer laser) which can completely change surface parameters. The different aspects are presented in two steps: the first consists in the surface characterization of the composite material and the second is related to results of destructive single lap shear tests of composite assemblies. Finally, both steps are linked in order to derive general rules on phenomena governing adhesion properties of polymer composites.


20.2 Materials and Methods 20.2.1 Composite Materials

Carbon/epoxy and glass/epoxy composites are manufactured by an autoclave process. The epoxy thermoset resin remains exactly the same for both composites; only the fiber reinforcement (glass or carbon weaving) differs. Due to the high performance demanded for aeronautic applications, the average fiber volume of these composite materials is generally above 60%. Unfortunately, owing to the industrial application (Hurel Hispano, Snecma group) of these composites, a detailed description of chemical compositions or technical information on these materials is proprietary. 20.2.2 Surface Analyses

The contact angle of sessile drops was measured using a GBX goniometer (Digidrop++ model). Uniform 3 lL drops of each liquid (water, glycerol, formamide, diiodomethane) were carefully deposited on composite surfaces at room temperature et humidity using Teflon syringes (5121 TLC-B Teflon) with an internal diameter of 0.73 mm. For each liquid used, the characterization was attained within 10 or 20 droplets on 20 mm ´ 50 mm surfaces. Both left and right contact angles were automatically calculated with the software (Windrop ++). The roughness was assessed by means of a laser interferometer apparatus (Altisurf 500®). Among all the available information deduced from roughness analyses, the parameter Sa determined from surface reconstitution [Eq. 1, where z is the hight difference of a given point having x, y coordinates], was used instead of the classical Ra. Indeed, Sa seems to be more representative of such rough anisotropic surfaces. Sa

X1 X1 N 1 M jzx; yj M N x0 y0

1

A surface morphology assessment and a study of the failure were both performed with a Hitachi S 3000 N scanning electron microscope (SEM) without surface metallization prior to observation, at a magnification ranging from 30 ´ to 2500 ´. Chemical analysis was carried out with SProbe (SSI) ESCA (electron spectroscopy for chemical analysis) apparatus. Analytical measurements were performed on an area of 300 lm ´ 1200 lm, at a depth of less than 10 nm.

307

308


20.2.3 Single Lap Shear Tests

Single lap shear tests were performed on 100 mm ´ 25 mm composite specimens with a 12.5 mm ´ 25 mm area bonded with an epoxy adhesive. Glass/ epoxy spacers were placed on one side of each bonded sample to assure perfect sample alignment during mechanical testing (Fig. 20.1). Shear stress measurements were carried out at room temperature at a displacement rate of 2 mm min–1 on an Instron 8802 apparatus with Merlin software. The stress necessary to break the joint was determined from the stress–time curves of five samples for each surface type, thus making it possible to calculate a representative average value and the typical error.

20.3 Results and Discussion 20.3.1 Why Excimer Laser Treatment?

Previous work performed on our composite materials [16] provided clear evidence of the influence of surface pollution prior to the use of any surface treatment on adhesion behavior. Several surface treatments can be performed to remove this surface pollution and increase single lap shear performances. As described in Section 20.1.2, peel ply treatment acts very specifically on composite materials. As several authors have already described [3, 5, 17] removing the peel ply before the application of the adhesive provides a rough (Fig. 20.2) and clean surface, free of pollution from the manufacturing process. Such surface cleaning, as well as several other surface modifications, leads to an increase in adhesion performance. Hence, failures which could previously occur adhesively between the composite surface and the adhesive layer will then propagate cohesively into the composite surface (Fig. 20.3) when this kind of treatment is applied, showing complete removal of the superficial polymer layer

Fig. 20.1 Single lap shear test samples.

Fig. 20.2 Surface profilometry of peel ply treated composite.


310

20 Laser Surface Treatment of Composite Materials to Enhance Adhesion Propertiess Fig. 20.3 SEM micrograph of a failed glass/epoxy sample. Shift of the failure mode from adhesive to cohesive in the material.

300 lm

when the assembly is broken. In that case, whatever the surface treatment used, no more adhesion enhancement can be achieved as the material itself becomes the weak point of the assembly. Alternatively, excimer laser treatment can be used to obtain additional improvement of adhesion. By this technique, polymer layers at the surface can be removed selectively before bonding, thus making the adhesive layer directly linked to the fiber reinforcement itself. Hence, crack propagation which occurs cohesively in the material under this polymer layer can no longer take place. Section 20.3.2 provides the results of such an approach. 20.3.2 Excimer Laser Surface Treatment 20.3.2.1 Surface Characterization Laser surface treatment can be used either below or beyond the ablation threshold of the surface. Laser treatment is more often used below the ablation threshold of the material, thus inducing efficient modification of the surface composition [12, 18]. Various laser parameters, such as the wavelength, the fluence (laser intensity), the nature of the environmental gas, or the pulse number, may be changed in order to modify the characteristics of the treated surface as the treatment induces the formation of polar chemical species (hydroxyls, carboxyls, peroxides, etc.). Therefore, the use of laser treatment below the ablation threshold induces adhesion improvement mainly through thermodynamic and chemical parameters. Furthermore, the laser treatment may be used beyond the ablation threshold of the polymer matrix. Then it induces partial or complete removal of the polymer, depending on the selected laser parameters. Hence, using a laser fluence above the ablation threshold of the polymer, but below that of the fiber reinforcement of a composite material, leads to complete removal of the polymeric layer without any fiber degradation [11]. As explained above, laser parameters are of primary importance in order to achieve suitable surface characteristics [19]. For example, as can be seen in Fig. 20.4, the increase in laser fluence gradually increases the ablation of the


a)

b1)

b2)

b3)

c1)

c2)

c3)

311

Fig. 20.4 Optical microscope images of the influence of laser parameters: (a) untreated; (b) fluence 150 mJ/cm2, (1) 25 pulses, (2) 100 pulses, (3) 250 pulses; (c) 100 pulses, fluence (1) 150 mJ/ cm2; (2) 240 mJ/cm2; (3) 500 mJ/cm2.

312


polymer matrix, thus bringing large areas of fiber reinforcement to the surface. Similar observations arise after increasing the pulse number, but this also increases the number of ablation cones which result from surface heterogeneities. The cones can be responsible for an improved mechanical interlocking of the adhesive into the adherend [19]. Hence, based on the parameters available, three different ablation rates were studied corresponding to a weak, a medium, and a full ablation according to laser profilometry (Fig. 20.5). The weakly ablated surface possesses the same surface roughness as an untreated sample (Sa = 0.6 lm). On the contrary, the medium and the fully ablated ones exhibit partial and total removal of the polymer layer at the surface of the composite, respectively. In the case of total ablation, the resulting roughness is directly related to the kind of fiber weaving used for reinforcement. It is important to note that even after high-fluence treatments, SEM makes it possible to obtain evidence of the absence of visible alterations to the fiber (not shown here). As the untreated surface and the weakly ablated one possess the same surface roughness parameter, contact angles may be compared. Hence for all the test liquids used, the contact angles are by 10–208 lower on the slightly ablated samples than on the untreated surfaces. Such a sharp increase in surface free energy after a slight laser treatment may arise for three different reasons: efficient cleaning of the composite surface; modification of the surface chemistry as the result of the initial ablation process [20]; or the growing influence of fiber reinforcement [21]. Actually, XPS analyses reveal high fluorine pollution (more than 32% of the surface atoms) for untreated samples. After 25 pulses of laser treatment, fluorine is no longer present at the composite surface. Moreover, a significant increase in the O/C ratio is observed simultaneously (0.14 : 1 and 0.21 : 1 for untreated and laser treated samples, respectively). Both the chemical analyses and the contact angle measurements indicate complete cleaning of the surface and/ or chemical modification brought by the polymer ablation. One can also observe that the whole composite surface composition is no longer affected when the number of pulses or the laser fluence is increased. Then, the O/C ratio remains at a value of 0.21 : 1.

20.3.2.2 Mechanical Tests In order to elucidate the influence of the modifications occurring on the composite surfaces, two adhesives were used. The first one (adhesive X) shows a poor adhesion with the composites used, and the second (adhesive Y) provides highperformance assemblies, with a predominantly cohesive failure in the composite at the initial stage. The mechanical behavior is quite different for the two adhesives (Fig. 20.6). Hence, the higher performance is not coming from the same surface treatment. In the case of adhesive X, the adhesive failure occurs principally for both untreated and peel ply treated surfaces. The corresponding improvement of lap

Fig. 20.5 Surface profilometry and surface roughness of glass/epoxy laser treated composites. (a) weak ablation (Sa = 0.6 lm).



Fig. 20.5 b Medium ablation (Sa = 13.1 lm).

314

Fig. 20.5 c Full ablation (Sa = 17.2 lm).


316


a)

b)

Fig. 20.6 Single lap shear results of glass/epoxy and carbon/ epoxy composites after several different laser treatments. (a) adhesive X; (b) adhesive Y.

shear strength is mainly related to roughening and partial cleaning of the surface after tearing off the peel ply. In contrast, weakly ablated samples show completely cohesive failure inside the adhesive, thus indicating that bonding between the adhesive and the polymer resin has been sharply enhanced. However, a further increase in ablation does not improve the mechanical performance of the resulting assembly. The picture changes more and more to adhesive failure

20.4 Conclusion

between the fiber reinforcement and the adhesive, thus illustrating the poor bonding between adhesive X and the fiber reinforcement. As described above, the use of laser treatment can be of primary interest for high-performance adhesives which show predominantly cohesive failure inside the material. To confirm this hypothesis, the use of the second adhesive (Y) indicates that for untreated, peel plied, and weakly ablated samples, no obvious change in lap shear performance is observed as all the assemblies break cohesively inside the material. Moreover, the medium ablation induces a sharp improvement in mechanical performance with the failure changing from cohesive inside the adhesive to cohesive inside the material. This result confirms that composite assemblies with high performance may be improved by bonding the composite material directly to its fiber reinforcement. Furthermore, one can observe that the fully ablated sample do not give the best performance. In this very special case, SEM micrographs clearly demonstrated that the failure occurs in a fiber layer near the composite surface. Laser ablation then induces weak boundary layers consisting of fibers weakly linked to the bulk material composite. In the light of these observations the interest of laser surface treatment is evident in terms of wettability, roughening, surface cleaning, and chemical modification. The use of two different adhesives makes it possible to check the relative effects of laser treatment on polymer modification as well as on fiber protrusion. When a cohesive failure inside the material is obtained after classical surface treatment, laser ablation makes it possible to improve the mechanical behavior by reaching the fiber reinforcement itself.

20.4 Conclusion

Adhesive bonding of anisotropic materials such as composites is complicated as it is governed by numerous parameters. As a key issue, the surface parameters prior to bonding have to be correlated with the mechanical performance of the resulting assemblies. It is clearly established that apart from any additional surface treatment, the manufacture of the composite material itself strongly affects its adhesion ability. Different surface treatments such as peel ply make it possible to roughen and clean composite surfaces and consequently to improve adhesion performance. However, like many other treatments, it is limited by the cohesive failure that may occur in the material. As an alternative, laser ablation appears to be quite an interesting surface pretreatment for polymer composites since it facilitates efficient control of the surfaces to be adhesively bonded. Furthermore, surface properties must be suitably defined by taking into account the nature of the composite material and the adhesive used.

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11

Chem., 1925, 88. A. F. Harris, A. Beevers, Int. J. Adhesion Adhesives 1999, 445–452. J. R. J. Wingfield, Int. J. Adhesion Adhesives 1993, 151–156. X. Roizard, M. Wery, J. Kirmann, Composite Struct. 2002, 223–228. L. W. Crane, C. L. Hamermesh, L. Maus, SAMPE J. 1976, 6–9. S. S.Voyutskii, Rubben Chem. Techn., 1957, 531. C. Asgeorges, L. Ye, M. Hou, Composites Part A 2001, 839–857. J. Comyn, L. Mascia, G. Xiao, Int. J. Adhesion Adhesives 1996, 301–304. J. F. Friedrich, S. Geng, W. Unger, A. Lippitz, J. Erdmann, Surf. Coatings Technol. 1995, 664–669. G. Kuhn, I. Retzko, A. Lippitz, W. Unger, J. Friedrich, Surf. Coatings Technol. 2001, 494–500. M. Rotel, J. Zahavi, S. Tamir, A. Buchman, H. Dodiuk, Appl. Surf. Sci. 2000, 610–616.

12 P. Laurens, M. Ould Bouali, F. Meducin,

B. Sadras, Appl. Surf. Sci. 2000, 211–216. 13 S. K. Koh, J.S. Cho, S. S. Yom, Y. W.

Beah, Curr. Appl. Phys. 2001, 133–138. 14 S. L. Nesbitt, J. A. Emerson, J. P. Bell, Int.

J. Adhesion Adhesives 2000, 429–436. 15 M. Tanoglu, S. H. McKnight, G. R. Pal-

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mese, J. W. Gillepsie Jr, Int. J. Adhesion Adhesives 1998, 431–434. Q. Bénard, M. Fois, M. Grisel, IUPAC 2004 World Polymer Congress Proc. 2004, 142. P. Molitor, V. Barron, T. Young, Int. J. Adhesion Adhesives 2001, 129–136. H. Hiraoka, S. Lazare, Appl. Surf. Sci. 1990, 264–271. H. C. Man, M. Li, T. M. Yue, Int. J. Adhesion Adhesives 1998, 151–157. Y. Novis, R. De Meulemeester, M. Chtab, J. J. Pireaux, R. Caudano, Br. Polym. J. 1989, 147–153. Q. Bénard, M. Fois, M. Grisel, P. Laurens, Euradh’04 Proc. 2004, 132–137.

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21 Effects of the Interphase on the Mechanical Behavior of Thin Adhesive Films – a Modeling Approach S. Diebels, H. Steeb, and W. Possart

Abstract

In the present conribution, we develop a continuum-based model to describe experimentally observable interphases in thin adhesive films. The model is based on an extended contiuum theory, i.e. the mechanical behaviour in these interphases is captured by an additional field equation. The introduced scalar order parameter models the microscopical mechanical properties of the film phenomenologically. As the film properties at the boundary to the substrate are thought to be extrinsic quantities, they are prescribed as Dirichlet-type boundary conditions in the present model.

21.1 Introduction

Thin polymer films are used as adhesives in bonding technology. The number of applications has been steadily increasing during recent years. Therefore, it has become necessary to improve the existing models which predict the mechanical properties of these films. The bulk behavior of polymers can be characterized in standard mechanical tests and this, on the one hand, motivates the formulation of classical continuum mechanical models including viscosity, plasticity, and damage on a macroscopic scale, besides the well-known nonlinear elasticity models (for an overview see, e.g., Ref. [31]). On the other hand, micromechanical models based on maximum entropy principles have been established taking into account the individual properties of the polymer network, i.e., the entanglement density, chain stretching, etc. [19, 42–44]. By applying homogenization procedures on the microscale effective properties can be transformed to the macroscale, and therefore macroscopic constitutive equations can be formulated on the basis of this upscaling procedure. Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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21 Effects of the Interphase on the Mechanical Behavior of Thin Adhesive Films

While for the bulk material the microstructural behavior is understood quite well, the situation becomes worse for small specimens, e.g., thin polymeric films. With thin films the aspect ratio h=k is the ratio of the specimen size (here the thickness h of the film) with respect to a characteristic microscopical length scale k (e.g., the characteristic length of the chain molecules). One reason for this is found in the existence of interphases in the polymer. Various investigations have shown that relatively thick interphases exist close to the boundaries of the polymer film [6, 7]; these interphases become more and more important if the thickness of the film decreases. In the interphases nearly every property changes with respect to the bulk properties, including chemical composition and entanglement density. A mechanical consequence was shown in Refs. [24] and [37]. It was found that the mechanical property, e.g., the local stiffness, also changes in the interphase. Furthermore, the change of the properties in the interphase depends not only on the polymer itself but also on the properties of the substrate. For instance, the properties in the boundary layer close to a polymer–glass interphase differ from the properties of a polymer–aluminum interphase [24, 37]. With respect to the mechanical response, the existence of the interphases leads to a dependence of the effective mechanical moduli of the polymer film on the film thickness and on the combination of polymer and substrate. Thus, thin films are expected to show a pronounced size effect. Due to the complexity of the formation of interphases, a completely satisfying microscopic interpretation of these effects cannot be given today, especially since the process of the interphase formation is not yet understood in detail. Therefore, a micromechanical model cannot be devised for calculating the global effective properties of a thin polymer film including the above-mentioned size effects governed by the interphases. On the other hand, a classical continuum-based model is not able to include any kind of size effect. An alternative to the above-mentioned classical continuum or the microscopical model is the formulation of an extended continuum mechanical model which, on the one hand, makes it possible to capture the size effect but, on the other hand, does not need all the complex details of the underlying microstructure of the polymer network. The formulation of extended continua can be traced back to the Cosserat brothers [13], who enriched the standard continuum by rotational degrees of freedom. As a consequence, such a continuum does not only transfer stresses but also couples stresses. The Cosserat or micropolar theory is well established now [17, 22, 32]. It is successfully applied if materials with microstructure are modeled on the macroscale, e.g., the equivalence of the Cosserat model is shown for materials with lattice-like microstructures [1, 2, 16, 25, 34, 35, 41] and with granular materials [3–5, 8, 15]. Furthermore, the micropolar approach is taken into account if strain localization arises which in the framework of the standard continuum would lead to a mesh dependence in the numerical solution (see, e.g., Ref. [29]). The Cosserat model introduces a so-called internal length k in the theory which allows for the existence of solutions of the bound-

21.1 Introduction

ary layer type and, as a consequence, these models capture size effects. From a theoretical point of view, the size effects can be used to quantify the additional material parameters [26, 27]. Beside the Cosserat model there exist further theoretical approaches such as gradient enhanced continua or nonlocal continua (see, e.g., Refs. [9, 17, 18]) which implicitly take care of a microstructure, and therefore, make it possible to predict size effects. Even if so-called kinematically extended models such as micropolar or nonlocal continua allow for a theoretical description of size effects, they are not very suitable in the present case. The main problem is that continua of this type provide size effects only of the type “smaller is stiffer”, which means that the stiffness does not decrease due to a reduction of the geometry. This results from the fact that physically meaningful boundary conditions lead to additional constraints and, therefore, to an increasing stiffness in the boundary layer. For beam- or lattice-like microstructures this is quite obvious, as the effective stiffness of the specimen is greater if the total rotation of the beams are fixed at the boundary. An alternative approach to such kinematically extended continua is presented by Capriz et al. [9–11] and Svendsen [39, 40]. They introduce scalar-valued order parameters to describe the influence of the microstructure in analogy to the Ginzburg-Landau theory. The starting point is an extension of the energy of the system under study. It is assumed that the additional microstructural parameters contribute to the kinetic energy, to the energy flux, and to the energy supply of the system. In order to obtain additional balance equations for the scalar-valued parameters, invariance principles are applied to the balance of energy. Furthermore, the entropy principle is taken into account. The resulting equations are of the form of the so-called balance of equilibrated forces derived by Goodman and Cowin [20, 33] or Passman et al. [36] in order to describe the behavior of elastic materials with voids or granular media. The advantage of this approach over kinematically extended continua is twofold: first, the additional parameter is abstract and does not require special assumptions of the local kinematics as it is necessary in micropolar continua, for example. As a consequence, the corresponding boundary conditions can be formulated more flexibly. Second, the coupling between the macroscopic balance of momentum and the additional balance of equilibrated forces is obtained constitutively, via the free Helmholtz energy function and not via the free energy function in combination with measures of deformation assumed a priori. Therefore, such an approach allows for size effects in both ways – either “smaller is stiffer” or “smaller is weaker” depending on the boundary conditions. In order to take into account the formation of the interphases in thin films, the present approach is based on an extension of the energy functional. The theoretical framework is briefly discussed in Section 21.2. A simple form of the free energy is proposed which allows for both types of size effects, depending on an equilibrium value of the additional order parameter and on the corre-

321

322


sponding boundary conditions. While the equilibrium value is used to describe the bulk properties, the boundary data may be used to characterize the combination of the polymer and the substrate.

21.2 Theoretical Framework

The starting point of the investigation is the introduction of a scalar microstructural parameter j which contributes to the total energy E of the body under study as pointed out in Refs. [38] and 39]. In Eq. (1) q, e and x_ are the mass density, the specific internal energy density and the velocity, respectively. The parameter k in the product q k describes microstructural properties and transfers the square of the rate of j to the dimensions of a specific energy density. In addition, the energy supply R1 and the energy flux R2 are also modified in the form of Eqs. (2) and (3), wherein q b is the body force density, q g is the supply of j, and q r is the heat supply. Further quantities are the stress vector t T n associated to the Cauchy stress tensor T and to the outer normal n, the microstructural flux s S n and the heat flux q q n. Z E

qe B

1 1 q x_ x_ q k j_ 2 dv 2 2

1

Z q b x_ q g j_ q r dv

R1

2

B

and Z t x_ s j_ qda

R2

3

@B

In this case the balance of energy may be written in the global format of Eq. (4), which may be localized by the standard argumentation yielding Eq. (5), where div is the divergence operator. dE R1 R2 ; dt

4

d 1 1 1 1 _ 2 q e q x_ x_ q k j _ 2 div x_ q e q x_ x_ q k j dt 2 2 2 2 divTT x_ S j_

q q b x_ q g j_ q r

5

21.2 Theoretical Framework

The classical balances such as the balance of mass [Eq. (6)] q_ q div x_ 0

6

and the balance of momentum [Eq. (7)] q x div T q b

7

can be obtained by invariance principles from the balance of energy [21]. The material time derivative d =dt is abbreviated to a superimposed dot. Furthermore, the balance of internal energy [Eq. (8)] can be derived from Eq. (5) by use of Eqs. (6) and (7). q e_

1 _ T : D divS j_ _ 2 q kj_ j q k j 2

q q g j_ q r

8

In Eq. (8), D is the rate of deformation and T : D is the stress power. In addition to the standard formulation of the balance of internal energy, some terms depending on j_ are present, which represent the energetic contributions associated with the microstructural parameter j. Note that it is not necessary to specify the physical meaning of the microstructural quantities j, S, and q k in terms of the underlying microstructure even if this is desirable in view of the interpretation. In order to obtain an additional balance equation for the microstructural parameter j, the principle of dissipation is utilized. The starting point is the entropy balance [Eq. (9)] with entropy density q g, entropy flux 'g , entropy supply rg, and entropy production ^g 0. d q g q g div x_ div 'g rg ^g dt

9

Following the usual arguments, the entropy flux and entropy supply are chosen to be given by Eq. (10), where H is the absolute temperature. Introducing the specific free Helmholtz energy density [Eq. (11)] transforms the balance of entropy into the Clausius-Duhem inequality [Eq. (12)], which is evaluated following the concept of Coleman and Noll [12]. Note that the results of the ColemanNoll procedure may be in general more restrictive than the results obtained by the Müller-Liu approach [28, 30]. Nevertheless, for the present application it is sufficient in order to describe the effect of the interphases observed in thin polymer films. Details of the evaluation process are given by Steeb and Diebels [38]. A discussion based on the Müller-Liu approach is given in Ref. [39]. In the present contribution only the basic results of the evaluation are summarized. 'g

q ; H

rg

qr H

10

323

324


We

gH

_ qgH

q e_

11 _ qW

div q H q grad

1 H

qr 0

12

While we do not want to give a sophisticated model including all the effects found in the mechanical behavior of polymers, we restrict ourselves to the simplest case, namely to an elastic small-strain model at constant temperature. Therefore, the governing variables are the linear strain tensor [Eq. (13)] derived from the spatial gradient of the displacement field u, and the microstructural parameter j and its gradient. The free energy density is assumed to be a function of the form of Eq. (14). e

1 grad u gradT u 2

W We; j; grad j

13 14

As discussed in Ref. [39] the following results [Eqs. (15)–(17)] are obtained from the evaluation of the Clausius-Duhem inequality if additionally q k is assumed to be constant: · Stress Tq

@W @e

15

· Flux of j Sq

@W @grad j

16

· Balance equation for the microstructural parameter qkj

qg q

@W @j

div S 0

17

Note that according to Eqs. (15) and (16) the free Helmholtz energy serves as a potential for the stresses T and for the microstructural flux S. According to the assumption of elastic material behavior, results of this type have to be expected. The additional balance equation for j [Eq. (17)] possesses the same structure as the balance of equilibrated forces obtained in Refs. [14, 20, 33] and applied, e.g., in Ref. [36]. Following the Müller-Liu approach, Svendsen [39] also derived a generalization of Eq. (17) for a model with scalar-valued structural parameters. In a last step the model is completed by an appropriate choice of the free energy function. The basic idea is to enhance the free energy of a linear elastic material [Eq. (18)] in such a way that the variation of the effective stiffness on the structural parameter is obtained, i.e., we assume that the Lamé constants depend on j in the form of Eq. (19).

21.3 Applications and Examples

W e l0 e : e

1 k0 tr e2 2

lj l0 expcj

j0 ;

18

kj k0 expyj

j0

19

The index 0 belongs to the bulk material, which is characterized by j0. If j j0 , then the bulk properties l0 and k0 are obtained. In Eq. (19) two additional material parameters c and y are introduced. As can be seen in Eqs. (16) and (17), derivatives of the free energy with respect to j and to grad j are also present in the governing equations. Therefore, additional contributions to the free energy with material parameters a and b are taken into account. For convenience, they are chosen to be of the quadratic type. Finally, the free energy density is chosen according to Eq. (20). W lj e : e kj tr e2

1 aj 2

j0 2

1 b grad j grad j: 2

20

This choice yields Eqs. (21) and (22) according to Eqs. (15) and (16) for the stresses T and the microstructural flux S, respectively. T 2 lj e kj tr e I

21

S b grad j

22

Furthermore, the abbreviation given by Eq. (23) is introduced, which may be interpreted as a configurational pressure [23]. p : q

@W a j @j

j0 lj c e : e kj y tr e2

23

21.3 Applications and Examples

A quasi-static version of the model is applied to describe the behavior of thin films in tension and shear experiments. The governing equations [Eqs. (24) and (25)] are obtained from the balance of momentum [Eq. (7)] and from the additional balance [Eq. (17)] if second-order time derivatives are neglected. div T q b 0

24

div S q g p

25

These equations are combined with the constitutive relationships for the stresses [Eq. (21)], for the microstructural flux [Eq. (22)] and for the configurational

325

326


Fig. 21.1 Q1P1 formulation with bilinear shape functions for u and for j.

pressure [Eq. (23)]. The solution to these equations is obtained numerically from a two-field finite element formulation. Details concerning the discretization are given in Ref. [38]. It was found that a Q1P1 formulation for the mixed functional leads to reliable results. According to Fig. 21.1, bilinear shape functions are chosen. 21.3.1 Uniaxial Tension Test

The first application of the model is a displacement-controlled uniaxial tension test. The geometry and the loading conditions are shown in Fig. 21.2. The values of the material parameters belong to a virtual material and are listed in Table 21.1. They are chosen in such a way that the effects become clearly visible.

Fig. 21.2 Geometry and loading conditions in uniaxial tension.

21.3 Applications and Examples Table 21.1 Material parameters.

Tension Shear

l0 [MPa]

k0 [MPa]

j0 [–]

c [–]

y [–]

a [MPa]

b [MPa m2 ]

10 000.0 10 000.0

5 000.0 5 000.0

1.0 1.0

1.0 1.0

1.0 1.0

5.0 200.0

10 000.0 10 000.0

Note that in addition to the standard Dirichlet boundary data for the displacements, Dirichlet boundary data for the microstructural parameter also have to be prescribed. Experimentally [24, 37] the local mechanical properties in the interphase depend on both the polymer and the substrate. The possibility of prescribing additional boundary conditions for j is utilized to describe the variations of the mechanical properties in the interphase depending on the substrate. j0 is chosen as Dirichlet data for the structural parameIf on the one hand j > j0 or ter, no interphase is predicted by the model. If on the other hand, j < j0 is chosen, an interphase is predicted which is either stiffer or weaker j than the bulk material, respectively. The thickness of the interphase is mainly governed by the material parameters a and b. =j0 2. As can be seen on the leftThe results shown in Fig. 21.3 belong to j hand side of the figure, the strain component e22 in the loading direction shows boundary layers while the stress component T22 in the direction of the applied load is constant. The strain in the boundary layers decreases compared with the inner part of the specimen, which means that the model predicts a stiff bound=j0 2. On the right-hand side of Fig. 21.3, the distributions of ary layer for j the microstructural flux S2 and of the microstructural parameter j itself are at the boundaries to j0 in the inshown. It can be seen that j decreases from j ner part of the specimen. If the same displacement is prescribed at the top of the specimen, but the Di=j0 0, a weak boundary layer is obtained. richlet data for j are changed to j This can be seen from the increasing strain component in the boundary layer as shown in Fig. 21.4. Again, the stress is constant in the specimen according to the equilibrium condition. The microstructural flux changes its direction and the parameter j now increases from the boundaries to the inner part of the specimen. The size effect is obtained by changing the length h of the sample while keeping the boundary data and the material parameters constant [38]. Due to the fact that the boundary layer is part of the solution and does not scale with the geometrical properties, its formation depends only on the boundary data and on the material parameters. Therefore, the thickness ratio of interphase to bulk material changes and, following simple mixing rules, the effective stiffness of the specimen as a whole changes. As a consequence, the effective Young’s modulus [Eq. (26)] depends on the thickness h. In Eq. (26) F T22 A is the resulting force, A is the cross-section and e Du=h is the effective strain according to the relative displacement of the load plates and the height of the specimen.

327

=j0 2. Fig. 21.3 Strain, stress, microstructural flux, and microstructural parameter in the specimen under uniaxial tension for j

328


=j0 0: Fig. 21.4 Strain, stress, microstructural flux, and microstructural parameter in the specimen under uniaxial tension for j 21.3 Applications and Examples 329

330


Eeff

F Ae

26

Note that even if the result Eeff Eeff h can be obtained from experiments, a theoretical approach will in general not be consistent if material parameters depend on geometrical data, such as the film thickness h. In the present example the situation is different: according to the extension of the theory by the parameter j, a consistent formulation is obtained. The material parameters are constant values but due to the enhancement the model is able to predict boundary layers as part of the solution. The introduction of the effective stiffness according to Eq. (26) is only a re-interpretation of the results, i.e., it is based on the solution of the boundary value problem. 21.3.2 Simple Shear Test

In the same way a shear test can be simulated. The geometry and the boundary conditions are shown in Fig. 21.5, while Figs. 21.6 and 21.7 show the distribution of the shear strain, the shear stress, the microstructural flux, and the microstructural parameter, respectively. Again, according to different Dirichlet data for j, both types of boundary layers are generated. The results are qualitatively the same as in the tension test.

21.4 Conclusion

In the present study an extended continuum mechanical model is derived which is able to predict either weak or stiff boundary layers in thin films. As a possible application, the formation of interphases in polymer films is investigated. In this case it was shown [7, 24, 37] that the local stiffness in the polymer depends on the combination of polymer and substrate.

Fig. 21.5 Geometry and loading conditions for the shear test.

=j0 2: Fig. 21.6 Field quantities for stiff boundary layers for j

21.4 Conclusion 331


=j0 0: Fig. 21.7 Field quantities for weak boundary layers for j

332

References

The enhancement of the model is based on the introduction of a microstructural parameter which contributes to the energy of the system. An additional balance equation of the type of the balance of equilibrated forces [20] is derived from the principle of dissipation in combination with the energy balance. The advantage of this procedure is the possibility of imposing boundary data for the structural parameter which capture the influence of the substrate on the formation of the interphases. An additional equilibrium value of the parameter may be given to characterize the bulk properties of the polymer. Therefore, giving appropriate values for the equilibrium quantities and for the boundary data, respectively, makes it possible to obtain the right distribution of the local stiffness in the specimen as a function of the structural parameter instead of geometrical data. In contrast to kinematically extended continua such as micropolar, micromorphic, or gradient continua, the chosen approach allows for weak and for stiff boundary layers as well. Additionally, in terms of an overall evaluation, the effective material parameters are obtained as a function of the layer thickness. This size effect can be utilized to determine additional material parameters included in the model. In further investigations, the material parameters have to be identified. This requires comprehensive mechanical tests on layers of different thickness. Furthermore, in order to describe the mechanical properties correctly, additional effects such as viscosity and/or plasticity have to be included into the formulation of the theory.

References 1 A. As¸kar, A.C ¸akmak (1968). A structural

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3

4

5

model of a micropolar continuum. Int. J. Engng. Sci., 6, 583–589. Z. Baz˘ant, M. Christensen (1972). ‘Analogy between micropolar continuum and grid frameworks under initial stress. Int. J. Solids Struct., 8, 327–346. M. Becker, W. Hauger (1982). Granular material – Experimental realization of a plastic Cosserat continuum? In Mechanics of Inelastic Media and Structures (eds.: O. Mahrenholtz, A. Sawczuk), pp. 23–39. M. Becker, H. Lippmann (1977). Plain plastic flow of granular model material. Experimental setup and results. Arch. Mech., 29, 829–846. D. Besdo (1985). Inelastic behaviour of plain frictionless block-systems described as Cosserat media. Arch. Mech., 37, 603– 619.

6 C. Bockenheimer (2003). Epoxid und Alu-

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minium im Kleberverbund nach mechanischer Vorbehandlung und Alterung. PhD thesis, Saarland University. C. Bockenheimer, V. Valeske, W. Possart (2002). Network structure in epoxy aluminium bonds after mechanical treatment. Int. J. Adhesion Adhesives, 22, 349– 356. N. Bogdanova-Bontcheva, H. Lippmann (1975). Rotationssymmetrisches ebenes Fließen eines granularen Modellmaterials. Acta Mech., 21, 93–113. G. Capriz (1980). Continua with Microstructures (Springer Tracts in Natural Philosophy, Vol. 35). Springer, New York. G. Capriz, P. Podio-Guidugli (1983). Structured continua from a Lagrangian point of view. Ann. Mat. Pura Appl., 135, 1–25. G. Capriz, P. Podio-Guidugli, W. Williams (1982). On balance equations for

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materials with affine structure. Meccanica, 17, 80–84. B. Coleman, W. Noll (1963). The thermodynamics of elastic materials with heat conduction and viscosity. Arch. Rat. Mech. Anal., 13, 167–178. E. Cosserat, F. Cosserat (1909). Théorie des corps déformables. A. Hermann, Paris. S. Cowin, J. Nunziato (1983). Linear elastic materials with voids. J. Elasticity, 13, 125–147. S. Diebels, W. Ehlers, T. Michelitsch (2001). Particle simulation as a microscopic approach to a Cosserat continuum. J. Phys. IV, 11, 5203–5210. S. Diebels, H. Steeb (2002). The size effect in foams and its theoretical and numerical investigation. Proc. R. Soc. Lond. A, 458, 2869–2883. C. Eringen (1999). Microcontinuum Field Theories, Vol. I: Foundations and Solids. Springer, Berlin. C. Eringen (2001). Microcontinuum Field Theories, Vol. II: Fluent Media. Springer, Berlin. U. W. Gedde (1995). Polymer physics. Chapman & Hall, London. M. Goodman, S. Cowin (1972). A continuum theory for granular materials. Arch. Rat. Mech. Anal., 44, 249–266. A. E. Green, R. S. Rivlin (1964). On Cauchy’s equations of motion. Z. Angew. Math. Phys., 15, 290–292. C. Kafadar, C. Eringen (1971). Micropolar media – I. The classical theory. Int. J. Engng. Sci., 9, 271–305. N. Kirchner, K. Hutter (2003). Modelling particles size segregation in granular mixtures. In Dynamic Response of Granular and Porous Materials under Large and Catastrophic Deformations (Eds.: K. Hutter, N. Kirchner). Springer-Verlag, Berlin. J. K. Krüger, W. Possart, R. Bactavatchalou, U. Müller, T. Britz, R. Santuary, P. Alnot (2004). Gradient of the mechanical modulus in glass–epoxy–metal joints as measured by Brillouin microscopy. J. Adhesion, 80, 585–599. R. S. Kumar, D. L. McDowell (2004). Generalized continuum modeling of 2-D periodic cellular solids. Int. J. Solids Structures, 41, 7399–7422.

26 R. Lakes (1986). Experimental microelas-

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ticity of two porous solids. Int. J. Solids Structures, 22, 55–63. R. Lakes (1995). Experimental methods for study of Cosserat elastic solids and other generalized elastic continua. In Continuum Methods for Materials with Microstructures (Ed.: H. Mühlhaus), pp. 1– 25. John Wiley, Chichester. I.-S. Liu (1972). Method of Lagrangian multipliers for exploitation of the entropy principle. Arch. Rat. Mech. Anal., 46, 131–148. H. Mühlhaus (1995). Continuum Methods for Materials with Microstructures. John Wiley & Sons, Chichester. I. Müller (1971). The coldness, a universal function in thermoelastic bodies. Arch. Rat. Mech. Anal., 41, 319–332. I. Müller, P. Strehlow (2004). Rubber and Rubber Balloons. Springer, Berlin. W. Nowacki (1986). Thermoelasticity. Pergamon Press, Oxford. J. Nunziato, S. Cowin (1979). A nonlinear theory of elastic materials with voids. Arch. Rat. Mech. Anal., 72, 175– 201. P. Onck (2002). Cosserat modeling of cellular solids. C.R. Mecanique, 330, 717– 722. P. Onck, E. Andrews, L. Gibson (2001). Size effects in ductile cellular solids. Part I: Modeling. Int. J. Mech. Sci., 43, 681– 699. S. Passman, J. Nunziato, E. Walsh (1984). A theory of multiphase mixtures. In Rational Thermodynamics (Ed.: C. Truesdell), pp. 286–325. Springer, Berlin. R. Sanctuary, R. Bactavatchalou, U. Müller, W. Possart, P. Alnot, J. Krüger (2003). Acoustic profilometry within polymers as performed by Brillouin microscopy. J. Physics D: Appl. Phys., 36, 2738–2742. H. Steeb, S. Diebels (2004). Modeling thin films applying an extended continuum theory based on a scalar-valued order parameter – Part I: Isothermal case. Int. J. Solids Structures, 41, 5071–5085. B. Svendsen (1999). On the thermodynamics of thermoelastic materials with additional scalar degrees of freedom. Continuum Mech. Therm., 4, 247–262.

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42 L. R. G. Treloar (1943). The elasticity of a

modeling of materials with kinematic structure. Acta Mech., 152, 49–80. 41 C. Tekoglu, P. Onck (2003). A comparison of discrete and Cosserat continuum analyses for cellular materials. In Cellular Metals and Metal Foaming Technology (Eds.: J. Banhart, M. Ashby, M. Fleck). MIT Verlag, Berlin.

network of long-chain molecules. I. Trans. Faraday Soc., 39, 36–41. 43 L. R. G. Treloar (1943). The elasticity of a network of long-chain molecules. II Trans. Faraday Soc., 39, 241–246. 44 L. R. G. Treloar (1946). The elasticity of a network of long-chain molecules. III. Trans. Faraday Soc., 42, 83–94.

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22 Effect of the Diblock Content on the Adhesive and Deformation Properties of PSAs Based on Styrenic Block Copolymers C. Creton, A. Roos, and A. Chiche

22.1 Introduction

Pressure-sensitive adhesives (PSAs) are increasingly used for demanding applications requiring a precise tuning of the properties of the adhesive. Yet the level of understanding of the relationship between the chemical structure of the adhesive (controlled by the synthesis tools available to the synthetic chemist) and the final application properties remain far from complete, and optimization of PSAs is still mainly done empirically. Several authors have pointed out that the practical work to detach a PSA from a solid surface (if the debonding is achieved by peeling a tape, this energy per unit area corresponds to the peel force per unit width of the tape) is highly dependent on the rheological properties of the adhesive [1–5]. This conclusion mainly stems from the experimental observation that “adhesion” master curves can be constructed by using the same time–temperature shift factors as are used for linear viscoelastic properties. However, this description remains to this day rather qualitative, since it is not at all obvious exactly which rheological properties are related to the adhesive energy, how to quantitatively predict the peel force from rheological properties, or what is the role played by surface energetics in the final peel force that is measured. Images of the peeling of a PSA from a rigid substrate [5–10] invariably show the formation of filaments joining the two surfaces and sometimes of cavitation preceding the formation of these filaments [8, 10]. It is clear that the material is highly deformed in these filaments, i.e., in a regime where linear elasticity or linear viscoelasticity no longer holds. The mechanisms of debonding of viscoelastic materials from a solid surface when adhesion is simply due to van der Waals forces should be the starting point of our analysis. Two types of mechanism are typically observed: either the viscoelastic material has a relatively high elastic modulus of the order of 1 MPa or above and one typically observes propagation of a crack as described schemaAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

338

22 Effect of the Diblock Content on the Adhesive and Deformation Properties of PSAs

tically in Fig. 22.1, or the viscoelastic material has a modulus well below 1 MPa and in this case one typically observes that the final separation is preceded by the formation of fingers and cavities leading to filaments as shown in Fig. 22.2. The first situation has been studied both experimentally [11–14] and theoretically [15–21] within the framework of linear viscoelasticity, i.e., in a situation where the viscoelastic materials remain in the linear regime. Although details and geometries differ, when a crack propagates in a viscoelastic material the strain rate that the material experiences varies spatially: it is very high near the crack tip and much lower far away from it. Since the linear viscoelastic properties of the material depend on strain rate, the value of G' and G'' vary spatially. For the simple case where the viscoelastic material has a relaxed modulus G? and an unrelaxed modulus G0 and a single relaxation time s, the propagating crack has the characteristic trumpet shape first proposed by de Gennes [22, 23] and shown schematically in Fig. 22.1. For this simple case, an expression directly relating the velocity-dependent fracture toughness Gc to the dynamic viscoelastic moduli of the adhesive can be obtained [19, 23]. A limitation of the model is that it assumes that the zone where nonlinear elasticity is important is small relative to the sample size. While this may be true when a crosslinked elastomer is debonded from a solid surface, the scanning electron micrograph of Fig. 22.2 shows that this is not generally the case for PSAs. Therefore, when investigating the effect of a change in molecular structure or formulation on the adhesive properties of a PSA it is not sufficient to characterize the linear viscoelastic properties of the PSA but it is also important to measure the nonlinear elastic properties of these materials.

Fig. 22.1 Schematic of the propagation of a crack at the interface between a hard surface and a soft adhesive.

22.2 Block Copolymer Based Adhesives Fig. 22.2 Edge view by scanning electron microscopy of a T-peel test done on a block copolymer based PSA. Photo courtesy of Ken Lewtas (Exxon Mobil Chemical).

Back in the 1960s Kaelble studied very thoroughly the peel process of PSAs from a hard surface [1, 2, 24]. One of his important findings was that the stress under the fibrillar zone preceding total debonding was not homogeneous but had a compression area, followed by a tensile peak and a plateau. Such a curve is reproduced from the original reference in Fig. 22.3 a. A spatial measurement of the stress in an experiment on the steady-state peeling of a PSA is not trivial to set up. However, the flat probe test is able to reproduce closely the stress distribution of the peel test, no longer as a function of spatial position but as a function of time. The analogy, as also noted by Chuang et al. [25], is evident by comparing Fig. 22.3 a and b. In a typical probe test, the probe is brought in contact with the surface of the adhesive at a velocity Va and subsequently removed at a constant velocity Vdeb. The stress applied to the adhesive film is at first compressive, then becomes tensile until failure is initiated, typically by the formation of cavities. The plateau stress is then representative of the high-strain deformation of the cavity walls in the tensile direction. This geometry is ideally suited to characterize the adhesive properties of a fibrillating PSA and in particular of two specific properties, namely its resistance to the formation of cavities under a nearly hydrostatic stress, and its ability to form stable fibrils upon debonding [26].

22.2 Block Copolymer Based Adhesives

Among the commercially available families of PSA one distinguishes three large groups: solvent-borne acrylic adhesives; emulsion acrylic adhesives; and block copolymer based adhesives [27]. The block copolymer based adhesives are typically formulated as a blend of one or several block copolymers and one or several tackifying resins. The block copolymers are based on a glassy monomer such as styrene and an elastomeric monomer such as isoprene or butadiene. The glassy sequence is the minority phase (15–30 wt.%) and diblock and tri-

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22 Effect of the Diblock Content on the Adhesive and Deformation Properties of PSAs Fig. 22.3 Schematic of a peel test (after Ref. [2]) showing (a) the stress distribution as a function of position under the peel front and (b) schematic of a force–time curve obtained from a probe test.

block architectures are typically used as shown schematically in Fig. 22.4. These block copolymers are interesting in practice because adhesive films can be processed with the hot-melt technique. At ambient temperature, the glassy phase forms small nodules which act as a filler and as physical crosslinks for the elastomeric matrix. However, above the critical order–disorder temperature the material becomes disordered and the physical crosslinks disappear [28]. If the mo-


Fig. 22.4 Schematic of the domain organization of the block copolymer chains and the tackifying resin in the blend. Both triblock copolymers and diblock copolymers are represented.

lecular weight is not too high, which is typically the case, the copolymer becomes a fluid that is easily deposited as a film on a substrate without the help of an additional solvent. Pure copolymers used as base materials for PSA applications can show some self-tack (by interdiffusion) but are typically not tacky at all on solid surfaces. In order to obtain PSA properties, the entanglements of the rubbery phase must be diluted to lower the elastic modulus of the physically crosslinked gel, and the material needs to be more dissipative to resist crack propagation at the interface [29]. Both modifications of properties are achieved through formulation with a low molecular weight but high-Tg tackifying resin which is typically miscible with the rubbery phase but immiscible with the glassy domains. Here we will focus on a typical block copolymer based PSA system where the glassy block is made of polystyrene (PS) and the rubbery one of polyisoprene (PI). The effect of the tackifying resin on the rheological and adhesive properties is well documented [30–32] and amounts to a dilution of the entanglement or crosslink density combined with an adjustment of the glass transition temperature (and therefore of the range of usage temperatures) of the PSA. We will therefore only summarize here the main findings. Fig. 22.5 shows the linear viscoelastic properties, at a reference temperature of 20 8C, of a SIS triblock copolymer with and without tackifying resin. As one can readily see, the master curve displays the typical behavior of a crosslinked rubber for the pure block copolymer. When the resin is added (in this case 60 wt.% resin and 40 wt.% polymer), the plateau modulus shifts to lower values while the characteristic increase in modulus due to the glass transition shifts to lower frequencies, corresponding effectively to an increase in the Tg at fixed frequency. The crosslinked rubber behavior comes from the presence of styrene nodules which act as physical crosslinks. In the pure copolymers they can readily be seen by atomic force microscopy (AFM) (Fig. 22.6 a) in tapping mode or by small-angle scattering (SAXS) with the dis-

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Fig. 22.5 Master curves showing the effect of the addition of 60 wt.% of Escorez 5380 tackifying resin on G' (continuous lines) and G'' (broken lines) of the SIS block copolymer. Reference temperature: 20 8C.

a

b

Fig. 22.6 (a) AFM micrograph of pure SIS copolymer and (b) small-angle X-ray scattering spectra of the SIS copolymer with various amounts of tackifying resin. The intensity logarithm has been shifted for clarity.

tinct presence of a scattering peak representing the average domain distance (Fig. 22.6 b). Of course, the degree of long-range order which can be achieved in these systems will depend on annealing conditions. The samples shown in Fig. 22.6 have been slowly dried in air and annealed at 40 8C for 24 h under vacuum: ordered structures with much better long-range order can be achieved with longer annealing times at higher temperatures [33]. As shown in Fig. 22.6 b, one of the additional effects of adding the resin is a considerable re-

22.2 Block Copolymer Based Adhesives Table 22.1 Composition of the four adhesives studied here.

Polymers Tackifier

SIS SI

100 – 150

81 19 150

58 42 150

46 54 150

duction of the degree of long-range order, which is clearly observed by SAXS and which was also confirmed by mean field simulations [34]. The main scattering peak almost disappears when an amount of resin typical of adhesive applications is added to the blend, and the secondary peak vanishes for even lower resin contents. Yet, rheological properties confirm the presence of the physical crosslinks. The structure is therefore that of randomly dispersed styrene domains in an isoprene + resin matrix. This loss of long-range order is likely to be even more pronounced for the adhesives processed by the hot-melt technique, which are clearly out of equilibrium, and implies a reduction of the driving force for phase separation due to the presence of the resin. In practice, block copolymer based PSA are often formulated from base polymer blends of triblock and diblock copolymers in various proportions. Setting aside cost considerations, the reasons for using a certain blend or even a pure triblock copolymer are typically based on performance in standardized PSA tests such as loop tack, peel, or shear tests. Yet, the effects of adding diblocks to a triblock copolymer on the details of the mechanisms of debonding are not known. We worked with four model PSA blends; the molecular characteristics and the composition of the blends are summarized in Table 22.1. The base polymers were triblock copolymers and diblock copolymers generously provided by ExxonMobil Chemical and synthesized by Dexco. The tackifying resin was a hydrogenated C5 based resin commercialized under the trade name of Escorez 5380, also provided by ExxonMobil Chemical, with a Tg of 48 8C. This type of resin is essentially a low molecular weight polymer of a monomer containing stiff cyclic groups which cause the high Tg. It is designed to be miscible with the isoprene phase but not with the styrene. All adhesive blends contained 40 wt.% of base polymer and 60 wt.% of resin; within the base polymer the proportion of diblock was varied from 0 wt.% to 54 wt.%. In order to characterize the adhesive properties of the blends we used mainly the instrumented probe test developed in our own laboratory [35]. While it is in principle identical to the type of probe tester presented by Zosel [36] and currently commercialized by Stable MicroSystems under the name Texture Analyzer, our design has introduced some important differences, shown schematically in Fig. 22.7. We have added a 458 mirror and a video camera to be able to visualize the contact area through a transparent substrate, and a very stiff tripod to be able to adjust the parallelism between the adhesive film and the probe in order to maximize the contact area. Additionally, compared with the commercial instrument, the compliance of our probe tester is much lower, avoiding some of the interpretation problems associated with testing a stiff layer with a compliant

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Fig. 22.7 Schematic of the experimental design and setup used for the probe tests.


apparatus. More details on our experimental apparatus can be found elsewhere [35, 37]. The output of a probe test is a force and displacement measurement as a function of time, as shown schematically in Fig. 22.3 b. In order to easily compare curves obtained at different probe debonding velocities and for different layer thicknesses, it is customary to convert these curves in nominal stress vs. nominal strain curve by normalizing the force by the maximum contact area in the compression stage (often, but not always, the complete probe surface) and by normalizing the displacement by the initial thickness of the film. It is in this conversion stage that the correction for the compliance of the measuring apparatus is taken into account [37]. It is important to note that in the experimental results reported here, and unlike recently reported results on tack of liquids [38–41], the compliance of the apparatus does not influence the measured value of the peak stress: our materials are very solid-like and the compliance of the apparatus is never much greater than that of the layer. Nevertheless, since the compliance of the apparatus is of the same order of magnitude as that of the layer before cavitation occurs, a correction needs to be applied to the raw data to obtain the correct strain data for the adhesive. The question of the interpretation of the stress–strain curve needs to be addressed as well. The simultaneous video acquisition of the data makes it possible to interpret the general features of the curve in terms of specific deformation mechanisms. Fig. 22.8 shows the stress–strain curve obtained with a 120 lm thick SIS + resin blend adhesive layer, debonded at 10 lm s–1 from a smooth steel surface. The first peak in stress is due to the nucleation of cavities at the interface between the probe and the layer. These cavities nucleate on preexisting defects due to the mismatch in surface topography between probe and layer [42, 43]. If the pre-existing defects are large compared with a characteristic length scale given by the ratio of the surface tension of the adhesive and the low-strain elastic modulus, the growth of these cavities occurs at a level of stress which is independent of defect size and only depends on the deformation properties of the material [43, 44]. This stress level corresponds to the beginning of the fibrillation plateau, which is defined as rbf (see Fig. 22.8). On the other hand, if the contact between the probe and the adhesive only leaves small contact defects at the interface, a characteristic overshoot in stress due to the surface tension of the adhesive is observed, a sign of the growth of small defects, which become cavities in the bulk of the adhesive layer. This overshoot gives rise to a characteristic peak in stress (rmax) often interpreted as a measurement of adhesion but really only representative of the elastic properties of the adhesive coupled to the defect population. Characteristically, the value of the peak stress is not very dependent on the surface tension of the substrate [55]. Once the cavities are formed in sufficient numbers from the defects, the compliance of the layer increases dramatically and the force drops to a lower level before remaining roughly constant at a level which is independent of the initial population of defects. At this stage the layer is no longer confined and the test proceeds by the progressive extension of the cavity walls in the tensile direction.

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Fig. 22.8 Typical nominal stress–nominal strain curve, with corresponding images for an SIS + resin adhesive. Vdeb = 10 lm s–1.

The pressure inside the cavities is close to zero while the outside pressure is about 1 atm (1 bar). However because of the solid character of the material, the pressure difference is entirely taken by the cavities on the outside of the circular contact patch and the work done by the tensile force against the atmospheric pressure is negligible. Although it is tempting to interpret the plateau in stress after the maximum as a simple tensile test, there are some significant differences, the most important being that the strain level in the cavity walls is unknown and could well vary spatially. It is important to note that often these extended cavity walls are referred to as fibrils, a term implying independent columns. Fig. 22.8 clearly shows that for PSAs they are actually walls between cavities. The majority of studies focusing on the deformation of these walls have started from the assumption that the observation of a filament structure is characteristic of the behavior of a liquid and requires flow. In our case, Fig. 22.5 clearly shows that the material is a soft solid and cannot flow. These measurements were performed at low strains in the linear viscoelastic regime, however, and one could argue that at high strains flow is likely to occur, analogously to what is observed for yield stress fluids. In order to check this hypothesis, we performed an interesting experiment that can be done with a probe test setup [37, 45]: Instead of fully debonding the layer from the probe, the test can be stopped after the peak stress, and during the cavity wall extension stage. The motors are stopped and the stress is allowed to relax, giving an idea of the elastic or viscous nature of the adhesive material at high strains [45]. Results are


presented as curves of force as a function of time in Fig. 22.9 a, and the relaxation stage, normalized by the stress at the beginning of the relaxation, is highlighted in Fig. 22.9 b. After 2 min of relaxation, the stress has decreased by only about 25–30%, independently of extension. Although a detailed analysis of the wall relaxation process is under way, these results already clearly show the elastic nature of the adhesive foam formed upon debonding.

Fig. 22.9 Relaxation of the fibrils in a probe test for the SIS adhesive. (a) Force vs. time curves for consecutive tests where the displacement was stopped at various levels of deformation of the fibrils, and left to relax for 120 s. (b) Relaxation curves for different stops, normalized by the stress at the beginning of the stop.

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22.3 Effect of the Diblock Content on Adhesive and Deformation Properties

Generally PSAs are well known for their very viscoelastic behavior, which is necessary for them to function properly. It was therefore important to characterize first the effect of the presence of diblocks on the linear viscoelastic behavior. Since a comprehensive study on the effect of the triblock/diblock ratio on the linear viscoelastic properties of block copolymer blends has recently been reported [46], we characterized the linear viscoelastic properties of our PSA only at room temperature and down to frequencies of about 0.01 Hz. Within this frequency range all adhesives have a very similar behavior in terms of elasticity, as can be seen in Fig. 22.10. The differences appear at low frequency, a regime where the free isoprene end of the diblock chain is able to relax. This relaxation process is analogous to the relaxation of an arm of a star-like polymer [47], and causes G' to drop to a lower plateau modulus, the level of which is only controlled by the density of triblock chains actually bridging two styrene domains [46]. If we now turn to Fig. 22.11 a and b, which show probe test curves for the four adhesive blends at two different probe debonding velocities, the differences between the four adhesive blends are striking and although measured stresses are systematically lower for the low probe velocity, the plateau stress is clearly lowered by the presence of diblock while the maximum extension is increased. The question is, then, whether these differences can be predicted by the linear viscoelastic properties. When trying to relate the probe test curves, performed at a certain probe velocity, with the linear viscoelastic properties of the adhesives, measured at a certain pulsation, the question of the equivalence between probe velocity, Vdeb and frequency always occurs. In fact a probe test imposes a highly inhomogeneous deformation on an adhesive layer and no equivalence can be rigorously made between a non-steady state experiment, which deforms the material in a

Fig. 22.10 Elastic component of the complex modulus at 22 8C for the four adhesive blends.

22.3 Effect of the Diblock Content on Adhesive and Deformation Properties

wide range of strain rates which vary as the experiment proceeds, and a steadystate experiment such as that performed in a rheometer at a single pulsation value. However, it is useful at least to establish an approximate equivalence at least. In our experiments, the probe test curves show that, regardless of the adhesive and of the applied Vdeb, the initial portion of the plateau region always falls at e * 2, i.e., or equivalently at h/h0 = 3. This point on the curve corresponds to the situation where all cavities have expanded in the plane and the walls start to be extended in the tensile direction (see Fig. 22.8). Given that the initial adhesive layer thickness is 100 lm, the nearly uniaxial strain rate applied to the cavity wall at e = 2 is 0.005 s–1 and 0.5 s–1 for Vdeb = 1 lm s–1 tests and Vdeb = 100 lm s–1 tests, respectively. If one assumes that a strain rate of 0.5 corresponds to an equivalent frequency x/2p, significant differences due to the presence of diblock are expected at low strain rates, but not at high strain rates where almost all four blends have

Fig. 22.11 Probe test curves for four tackified SIS + SI blends at two different probe pullout rates. (a) Vdeb = 1 lm s–1; (b) Vdeb = 100 lm s–1.

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identical values of G'. According to Fig. 22.11 b, this is not the case: The curves remain very different at high strain rates.

22.4 Understanding the Structure of the Extended Foam

Given the high strains which are imposed on the adhesive layers in this wall extension regime, a systematic characterization of the nonlinear elastic properties of the adhesive is necessary. We performed uniaxial tensile tests of adhesive layers [48] up to high strains by casting the adhesive on release liners and then clamping the free-standing PSA films to a tensile testing machine fitted with a more sensitive load cell. The main results are summarized in Fig. 22.12 a and b, showing the nominal stress versus extension of the four adhesive blends at two different crosshead velocities chosen in order to be equivalent to the strain rate applied to the materials in the probe tests at the beginning of the plateau region. Clearly the differences in the nonlinear elastic properties of the adhesives, in the high-strain regime, are very pronounced, while the initial portion of the curve does not show very marked differences. However, in this representation a more quantitative comparison between the different adhesives is not straightforward. The stress– strain curves shown in Fig. 22.12 should be familiar to those readers working with rubbers; they display a fairly typical type of nonlinear elastic behavior observed for crosslinked rubbers. This similarity is logical since we have a structure of physically crosslinked styrene domains and a soft deformable isoprene + resin matrix. It is also very useful for the purpose of analyzing the data, since rubber elasticity has been a very active field for decades and several models, both phenomenologically and molecularly based, are available in the literature. The simplest model is the statistical theory of rubber-like elasticity, also called the affine model or neo-Hookean in the solids mechanics community. It predicts the nonlinear behavior at high strains of a rubber in uniaxial extension with Eq. (1), where rN is the nominal stress defined as F/A0, with F the tensile force and A0 the initial cross-section of the adhesive layer, k is the extension ratio, and G is the shear modulus. rN Gk

1=k2

1

Realizing the shortcomings of this simple model, extensive work has subsequently been performed to refine the picture of the nonlinear elastic behavior of crosslinked rubbers. Further information is available in textbooks [49] and recent reviews [50]. Among the various models that have been proposed to describe experiments, two will now be compared with our experimental data: a phenomenological model first proposed by Mooney [51] and further developed by Rivlin [52] which is based on the incompressibility condition and introduces a k-dependent term


Fig. 22.12 Tensile stress–strain curves for the four adhesives at two different strain rates. (a) Crosshead velocity 5 mm min–1 corresponding to an initial strain rate of 0.005 s–1. (b) Crosshead velocity 5 mm min–1 corresponding to an initial strain rate of 0.5 s–1.

in the modulus, and a recently proposed molecularly based model proposed by Rubinstein and Panyukov [50] which accounts very well for the respective role played by crosslinks and entanglements during the extension and compression of the rubber. In uniaxial extension, the Mooney-Rivlin model predicts Eq. (2) to hold, where C1 and C2 are two material constants. C2 rN 2 C 1 k k

1 k2

2

Note that, with this model, the reduced stress defined by Eq. (3) depends on the deformation k, whereas it did not for the simple affine model.

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rR k

rN 1 k2

3

If the curves of Fig. 22.12 a are plotted in terms of reduced stress as a function of 1/k, they appear as a set of parallel lines within the range of k–1 from 0.2 to 0.6 (Fig. 22.13). The slope of these lines gives the C2 constant directly while the intercept gives C1. It then becomes immediately apparent that the different adhesive blends have very similar values of C2 but very different values of C1. The upturn in reduced stress at high values of k is due of course to the strain hardening of the material, which is not captured by the Mooney-Rivlin model. The actual values of C1 and C2 obtained from fitting our experimental data are given in Table 22.2 for the four adhesive blends. The quality of the fit in the low to intermediate strain regime can be seen in Fig. 22.14, showing the experimental data and the fits using Eq. (2) and the parameters of Table 22.2. The relevance of Fig. 22.14 becomes obvious when comparing the curves of Fig. 22.14 with the portion after the peak of the probe test curves of Fig. 22.11. A striking feature of the probe test results is the large difference in the level of stress of the plateau region, which cannot be attributed to a variation in shear modulus G' or even in the complex modulus G* = (G'2 + G''2)1/2. On the other hand, these differences are well explained by the high-strain behavior of the adhesive in extension. Fig. 22.15 shows the ratio rbf/G* between the stress at the beginning of the plateau extracted from the probe test curves (at e = 2) and the complex low-strain modulus, and between rbf and the stress in uniaxial extension at k = 3 corresponding to the position e = 2 in the probe test (k = e + 1). Clearly the stress level in the plateau region of the probe test curve is directly related to the behavior of the adhesive in uniaxial extension at high strains,

Fig. 22.13 Reduced stress representation of the stress–strain curves for the four adhesives. The broken line is an illustrative fit of the data with the Mooney-Rivlin model.

22.4 Understanding the Structure of the Extended Foam Table 22.2 MooneyRivlin parameters [kPa] fitted from the uniaxial extension data.

C1 C2

0 wt% SI

19 wt% SI

42 wt% SI

54 wt% SI

17 44

9.5 45

2 50

0 54

which is quite different for the three adhesives, but is not related at all to the value of the shear modulus at low strains. This is one of the significant findings of our study, showing conclusively that the crosslinking network of the adhesive will be essential in controlling the stress level in the plateau and therefore the work done to detach the adhesive. From these encouraging results, we can attempt a more molecular interpretation and fit our uniaxial extension data to a molecularly based model such as that proposed by Rubinstein and Panyukov [50], which was developed to model crosslinked rubbers. Much effort has been dedicated to the development of better molecularly based models, which have been discussed in an excellent recent review [50]. Since our block copolymer blends are not strictly speaking rubbers (PS spheres occupy some volume and have a high functionality relative to chemical crosslink points), any quantitative comparison between our data and a molecularly based model should be taken with a bit of caution. We feel however that the insight provided by the molecularly based model is essential for the understanding of the mechanical properties of our systems. A particularly interesting molecular model for the nonlinear elasticity of polymer networks was proposed by Rubinstein and Panyukov [50], the slip-tube model. In this model the confining potential acting on the network chains depends on the deformation and is modeled by virtual chains attached to the network

Fig. 22.14 Experimental data (*) and best fit (full line) with the Mooney-Rivlin model in uniaxial extension.

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Fig. 22.15 Normalized values of the stress at the beginning of the plateau. Values of rbf/G* (s ) and of rbf/r (k = 3) (·).

chains through slip-links, allowed to slip along the contour of the tube but not to pass through each other. The slip-links represent the entanglement points, which are trapped by the crosslinking process. In the regime of interest, the nominal stress can be well approximated by Eq. (4), where Gc and Ge are two physical moduli given by Eq. (5), in which w is the network functionality, v is the total number of elastic strands per unit volume, c is the monomer concentration and Ne is the average number of monomers between entanglements. Therefore, Gc and Ge represent, respectively, the part of the elastic modulus due to the phantom network (fixed) and the part due to the entanglements (slip-links). Note that at low strains, i.e., for k?1, Gc + Ge is equivalent to the shear modulus G and for very high strains the stiffness of the material depends only on Gc, i.e., on the density of crosslinks. rN

k

Gc 1

1 Ge Gc k2 0:74 k 0:61 k 2 ckT vkT Ge w Ne

1=2

0:35

4 5

The parameters obtained from fitting Eq. (4) to the experimental data are shown in Fig. 22.16. The fits are not as good as the Mooney-Rivlin fits [45, 48] but nevertheless capture reasonably well the stress–strain curves in approximately the same ranges of extension ratios. The results can be interpreted as follows: the parameter Gc, which is directly related to the volume density of fixed crosslink points, varies significantly between the pure triblock adhesive and the high diblock content adhesives, where the fit gives a value close to zero. On the other hand the parameter Ge is much higher and nearly independent of the diblock content. This shows that the low-strain modulus is essentially con-


trolled by the entanglement network of the soft phase (isoprene + resin) and does not change much with increasing amounts of diblock, while the highstrain behavior is increasingly controlled by the apparent crosslink density, which decreases dramatically with increasing diblock content. A simple molecular picture would have the triblock chains forming bridges between styrene domains and providing the fixed crosslinks. However, a more detailed analysis shows that the variation in bridging triblock chains can only account for a small fraction of the decrease in effective crosslink points. It is likely that permanently trapped entanglements and slowly relaxing entanglements also provide a significant contribution to the fitted value of Gc and the role of the presence of diblock on this fraction of trapped entanglements is currently unknown. It is important to note that if quantitative comparisons between model and experiment are to be made, the filler effect of the styrene domains should also be taken into account in the prediction of the shear modulus [46, 53, 54]. However given the small volume fraction of styrene (of the order of 6%) in the blends, this effect will only add a factor of about 1.2 in the prediction of the modulus. In conclusion, when such an adhesive is debonded from a high energy surface such as steel, the high-strain properties of the adhesive control the formation and extension of the fibrillar structure which provides the bulk of the work necessary to detach the adhesive from the surface, and hence the major part of the peel force. We have seen that the level of the plateau stress can be predicted quantitatively by a simple tensile test. From the studies on cavitation, we know that the nominal stress at the plateau corresponds also to the cavity growth stress for large initial defects. Although the fibril extension stress can be predicted from the nonlinear elastic properties of the adhesive, in practice the important property that one wishes to predict is the adhesion energy rather than simply the plateau stress of the fibrillar zone. This prediction would require a better understanding of which molecular features control the detachment of the fibril from the surface, once it is

Fig. 22.16 Values of Gc (·) and Ge (s) obtained from fitting the experimental data to the slip-tube model.

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highly extended. This problem is currently still open since it involves not only knowing the nonlinear elastic properties of the material at very high strains (the strain hardening) but also a microscopic criterion for fibril detachment.

22.5 Interfacial Fracture

This brings us back to more interfacial properties. It is interesting to investigate what becomes of the effect of adding diblock in the adhesive when the PSA is detached not from steel but from a low-adhesion surface such as a polyolefin, a silicone layer, or any other release surface. We have carried out a comprehensive investigation of the debonding mechanisms of the adhesives of Table 22.1 from a surface of ethylene–propylene (EP) copolymer [45]. The copolymer layer was spin-coated from a xylene solution onto the flat end of the steel probe. Two representative probe test curves for the detachment of an SIS adhesive from steel and from EP surfaces are shown in Fig. 22.17: while the initial portion of the curve is identical, the force drops rapidly to zero for the EP surface, and never forms the characteristic fibrillar plateau observed on steel surfaces. How does this happen? As qualitatively described by Creton et al. [55] for a detachment from a polydimethylsiloxane layer, when the resistance to crack propagation Gc is low, cavities are nucleated (around the peak stress) and then propagate as interfacial cracks at the interface between the probe and the adhesive, and eventually coalesce. This process of crack propagation and coalescence is responsible for the sharp drop in force observed in Fig. 22.17 for the EP surface and occurs at rather low values of nominal strain. In this case no formation of the characteristic foam structure responsible for the high debonding energy is observed.

Fig. 22.17 Probe test curves of the adhesive containing 19 wt.% of diblock debonded at Vdeb = 1 lm s–1 on a steel surface and on an EP surface.


This balance between interfacial propagation and bulk deformation has been described for linear elastic materials [56] and results from the competition between two mechanisms: the velocity of propagation of an interfacial crack, which is controlled by the critical energy release rate Gc; and the bulk deformation, which is controlled by the cavitation stress and hence essentially by the elastic modulus E or G. In the linear elastic model, the key parameter is the ratio Gc/E, which represents the distance over which an elastic layer needs to be deformed before being fully detached from the hard surface. This model has been verified experimentally for elastic gels [57]. However, for PSA layers we need to introduce two modifications which complicate the analysis: the adhesives are both viscoelastic and strained in the nonlinear elastic regime. In other words the term Gc will include a dissipative term and the term E should be replaced with a high-strain equivalent controlled by the nonlinear elastic properties as shown in Fig. 22.14 and 22.15. As a result, the crack front will not have the same shape as the classical interfacial crack and the exact nature of the stress distribution at the crack tip will be unknown. A schematic description of the shape of the crack for the linear elastic case and for the PSA case is shown in Fig. 22.18 and although a discussion of the effect of high strains and viscoelasticity on the crack shape are beyond the scope of this paper, the interested reader is referred to the work of Hui and Jagota [58] or Newby et al. [59] on that topic. Results of probe tests with the four adhesive blends on the EP surfaces are quite different from the results obtained on the steel surfaces, as shown in Fig. 22.19. The high diblock content adhesives show a pronounced plateau, while the adhesives with low or no diblock content show only a shoulder after the peak. Although the magnitude of the effect depends on the probe velocity, the tendency is clear: the presence of diblock makes it easier to extend the cavity walls in the tensile direction. This effect is also illustrated by another experiment performed with our adhesives: the in-situ relaxation experiment [60]. In

Fig. 22.18 Schematic of a crack tip: the broken line represents the classical shape obtained from linear elasticity and the full line represents a possible evolution of the crack shape when finite deformations and non-

linear elasticity are taken into account. Note that the contact angle H is sensitive both to the mechanical properties of the adhesive and to the boundary conditions at the interface.

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Fig. 22.19 Probe test curves for the four adhesives debonded from an EP surface.

this type of experiment, a tensile stress clearly inferior to the measured peak stress is applied to the confined adhesive layer, in the probe test geometry, and the slow failure of the layer as a function of time is observed. During the fast loading portion of the test, some cavities nucleate (those growing from the larger defects); then the nominal displacement of the probe is stopped, and the system is allowed to relax. Because of the finite compliance of the measurement setup, the force relaxes and the thickness of the layer (i.e., the distance between the substrate and the probe) increases slowly, applying a slow deformation rate which is an intermediate situation between that of relaxation and that of creep. Observation of the layer through a microscope reveals which deformation mechanisms are responsible for this increase in average thickness. Since our adhesives are not fluids, most of the increase in thickness has to be taken up by the formation of voids and this is indeed observed experimentally [60]. However, the two series of images of Fig. 22.20 a and b reveal how these voids create free space. In Fig. 22.20 a, the adhesive is a pure triblock + resin adhesive: the sequence of images as a function of time reveals that no new cavity becomes optically visible during the force relaxation process, and the cavities formed during the fast loading step propagate as cracks at the interface between the probe and the film in a way reminiscent of a forced dewetting. On the other hand, in Fig. 22.20 b, the adhesive includes 42% SI in the formulation, and in this case existing cavities grow slowly and remain circular while new cavities become optically visible during the relaxation step. Since all the adhesives have the same monomer composition, this difference in behavior can only be attributed to the difference in rheological properties between the four adhesives. Although materials are highly viscoelastic, it is legitimate to make the analogy between the propagation of the crack front at the interface between the film and the probe and the propagation of a crack at the interface between a crosslinked rubber and a solid surface, which has been exten-


Fig. 22.20 Sequences of images from a slow debonding of two adhesives from a fluorinated surface: (a) pure SIS adhesive; (b) SIS + 42% SI adhesive.

sively studied [61, 62]. Resistance to interfacial crack propagation is related to the energy dissipated by the viscoelastic material: in steady-state conditions the energy released by the material plus the external work provided by the operator match the energy dissipated near the crack tip. However, the deformation field at the crack tip is complex and highly dependent on the precise shape of the crack front. If finite deformations are considered, the shape of the crack front can vary with propagation rate, material properties, and surface properties, as shown schematically in Fig. 22.18. Therefore the amount of energy dissipated at the crack tip for a given crack velocity is likely to be a complex function of the viscoelastic properties of the adhesive at low but also at high strains, as well as of the boundary conditions at the interface (slip, no slip, partial slip) [63–65]. The value of the dissipative factor tan d = G''/G' for the four adhesive blends is shown in Fig. 22.21. Clearly the addition of diblock in the blend has no effect on the dissipative properties of the adhesives at high frequencies but it has a significant effect at low frequencies. A relaxation experiment such as that described in Fig. 22.20 a and b involves a very slow growth of interfacial cracks, precisely in the regime where linear viscoelastic properties differ. We can therefore propose, at least qualitatively, that the spectacular improvement in adhesive properties observed for the high diblock adhesives on EP surfaces is due to their more dissipative character, which slows crack propagation considerably at the interface, therefore avoiding early coalescence between separate cavities, and favors the formation of a fibrillar structure with the cavity walls. An explanation such as this remains very simplistic, since a range of strains and strain rates coexist at the crack tip and it is unlikely that a property measured in the linear regime such as G' and G'' alone can account for the behavior

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Fig. 22.21 Dissipation factor tan d of the four adhesive blends at 22 8C (from pure SIS to 54 wt.% SI).

of the adhesive layer at the crack tip. An extensive study of the dissipative properties of the adhesive materials as a function of strain could reveal a more complex role for the presence of the diblock. Already the results of Fig. 22.14 show that the nonlinear elastic properties of the four adhesives are very different and their aptitude to form blunted cracks is going to be very dependent on their high-strain behavior [58].

22.6 Summary

Summarizing our results on the material properties, we observe that the presence of diblock copolymer in a block copolymer based adhesive formulation has several remarkable effects, some of which have never been reported and which affect the adhesive properties: · In the linear viscoelastic regime, the presence of diblock introduces an additional relaxation time which is apparent at low frequency and is due to the relaxation of the free end of the diblock in the elastomeric matrix. The more diblock in the adhesive, the more pronounced is the jump in modulus between the relaxed and the unrelaxed modulus. · The nonlinear elastic properties can be described by both the Mooney-Rivlin model and the molecularly based slip-tube model. Both of these models stress the fact that the low-strain modulus of the adhesives is controlled by the entanglement structure of the isoprene + resin phase, while the high-strain modulus is controlled by the physical crosslink structure. The incorporation of diblocks in the adhesive dramatically reduces the density of crosslinks and causes a more pronounced softening in the high-strain part of the stress– strain curve.

References

Both of these material properties have direct consequences on the adhesive properties: · On high-energy surfaces where a fibrillar structure can be formed, the cavitation stress for large defects and the level of plateau stress in the fibrillar regime are significantly higher for the pure triblock systems and are directly related to the density of physical crosslink points or trapped entanglements. The implication of this result is that pure triblock systems will perform best in PSA applications where initiation of failure is the limiting factor, such as long-term resistance to shear. However, for applications where a specific peel force is required, it is not clear that the pure SIS formulations will perform better since the maximum fibril extension is systematically reduced for the triblock based adhesives relative to the adhesives containing diblocks. · On low-adhesion surfaces, the addition of diblock should be much more beneficial since it provides two essential features which will slow down crack propagation: a more pronounced softening at intermediate strains and a more dissipative character. When failure is initiated by the formation of cavities, the softening behavior favors crack blunting and hence the formation of fibrils, while the more dissipative character slows down crack propagation and again makes it easier to form the foam structure. In terms of the parameter Gc/E, controlling the behavior of fully elastic systems, the more dissipative character increases Gc while the softening decreases the effective high-strain E. Hence for applications where the PSA is applied to low-adhesion surfaces such as polyolefins, a reasonable percentage of diblock should increase the practical work of adhesion.

Acknowledgments

We gratefully acknowledge the financial support of the European Commission under the GROWTH program of the 5th framework, Collaborative Project No. G5RD-CT2000-00202 DEFSAM. We are also indebted to Ken Lewtas, Jacques Lechat, and Galina Ourieva from ExxonMobil Chemical Europe for providing the materials and for helpful discussions.

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23 Contact Mechanics and Interfacial Fatigue Studies Between Thin Semicrystalline and Glassy Polymer Films R. L. McSwain, A. R. Markowitz, and K. R. Shull

Abstract

In order to develop a greater understanding of interfacial interactions between a semicrystalline polymer and a glassy polymer, adhesion tests were performed on very thin layers of poly(ethylene oxide) (PEO) sandwiched between two layers of poly(tetramethyl bisphenol A polycarbonate) (TMPC). The tests were designed to provide intimate contact between the surfaces while they were heated above the melting point of the PEO and cooled back to room temperature. A contact mechanics approach, based on the Johnson, Kendall, and Roberts (JKR) theory, was used to determine values of the energy release rate describing the energetic driving force for crack propagation within the interfacial region. The ability to measure crack propagation at large values of the energy release rate was limited by rupture of the silicone elastomer that was used to provide a sufficiently compliant matrix for the adhesion experiment. By cycling the tensile stress at relatively low loading levels, we were able to measure fatigue crack propagation at values of the energy release rate that did not result in failure of the elastomer.

23.1 Introduction

Although there has been significant research relating to the molecular origins of fracture and adhesion in glassy polymers [1], similar studies involving semicrystalline polymers have been much more limited. Examples where progress has been made include adhesion tests focusing on the study of copolymer formation between two immiscible semicrystalline polymers [2], the effect of grafted chains on semicrystalline adhesion [3, 4], and the increase in strength associated with cocrystallization of semicrystalline polymers at an interface [5]. Recent progress includes work by Plummer et al. that probed the use of models for glassy polymer fracture mechanisms to describe the failure associated with a semicrystalline system [6]. Laurens et al. have also asserted that the molecular Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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23 Contact Mechanics and Interfacial Fatigue Studies Between Semicrystalline and Polymer Films

mechanisms can be related to the crystalline orientation close to the interface [7]. These recent studies point to the fact that interactions between semicrystalline polymers are more difficult to characterize than those in glassy polymers because of the more complex microstructure of semicrystalline polymers [1]. Bidaux et al. focused on the in situ formation of a diblock copolymer at an interface between two immiscible semicrystalline polymers. In order to produce an adhesive interface, the samples were heated above the melting temperature of one of the polymers while they were in contact, so that mobility of the chains was possible. By heating up above the melting temperature of both of the polymers during bonding, the fracture energy increased up to 700 J m–2. This value was consistent with values for cohesive failure of the polymer with the lower melting temperature [8]. These very large values of fracture energy are associated with significant plastic deformation within the semicrystalline polymer. Benkoski et al. have utilized diblock copolymers, composed of a glassy block and a semicrystalline block, to reinforce an interface [9]. Their studies indicate that the penetration of the chains from the diblock into the homopolymer allow a transfer of stress across the interface that is dependent on parameters related to the crystalline chains in the diblock. By increasing the length of the crystalline portion of the diblock, values of the fracture energy increased from 1 to 700 J m–2 [9]. As with the experiments of Bidaux et al., the largest values of the fracture energy were attributed to plastic deformation and cohesive failure within the semicrystalline polymer. In order to develop a greater understanding of the interfacial aspects relating to fracture, tests can be performed on very thin layers of semicrystalline polymer that are sandwiched between two glassy polymer layers. By controlling the thickness of the semicrystalline layer, the volume of polymer that can be deformed is limited. Adhesion tests probing the effect of the thickness parameter of this middle semicrystalline layer will provide information on the stress transfer across the interface as well as the dissipation of energy within this layer. An appropriate adhesion test must be sensitive to the relatively low values of the fracture energy obtained in the absence of bulk energy dissipation mechanisms, and must enable intimate contact between the surfaces to be maintained while the temperature is cycled above and below the melting point of the semicrystalline polymer. In this paper, we approach this problem by utilizing an axisymmetric adhesion test involving the contact of a rigid glass hemispherical indenter with an elastomeric substrate. Application of these types of tests has been reviewed recently [10]. The concepts that are most important for understanding the data that is obtained are summarized in this section. The schematic in Fig. 23.1 illustrates the sample geometry and testing variables that are important for this test. The analysis, which was developed originally by Johnson, Kendall, and Roberts (JKR) [11], and later placed in a fracture mechanics context by Maugis and Barquins [12], can be used to obtain the energy release rate, G, which is the energetic driving force for crack propagation within the system [Eq. (1)]. In this equation, Pt is the applied tensile load, a is the contact radius, R is the radius of curvature of

23.1 Introduction Fig. 23.1 Sample geometry including a semicrystalline PEO layer (dark gray) sandwiched between glassy TMPC polymer layers (white).

the indenter, and K is the modulus of elasticity. For an incompressible material with Poisson’s ratio equal to 0.5, K = 16E/9, where E is Young’s modulus. G

Ka3 =R Pt 2 6p Ka3

1

The elastic modulus can be obtained independently from Eq. (2), which states the relationship between the load, displacement dt, and contact radius [12]. dt

a2 2Pt 3R 3Ka

2

The compliance, C, of the system can also be used to relate the radius of contact during the test to a change in the displacement [10, 13] according to Eq. (3). @dt 2 C @Pt a 3Ka

3

In our tests, a rigid glass hemispherical indenter and an elastomeric substrate, both coated with glassy polymer layers, are brought into contact. A thin layer of poly(ethylene oxide) (PEO) is sandwiched between the two glassy polymers, and serves as the adhesive layer. The interaction of the polymer layers is temperature-dependent, so the sample and indenter are heated while in contact. Once the samples have cooled back to room temperature, the elastomer is retracted from the indenter and the force of separation is measured. The experiment is conceptually similar to the approach utilized by Mangipudi et al., who used this approach to obtain the work of adhesion between glassy polymers at room temperature [14]. In these previous experiments the work of adhesion was dominated by equilibrium surface energetics, where the values of G are close to 0.1 J m–2. In our case we are interested in extending this methodology to the larger values of G characteristic of fracture processes in semicrystalline polymers. The compliance for these tests is dominated by the compliance of the silicone elastomer substrate; therefore, the fracture of the elastomer at high applied

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stresses limits the applicability of the test to values of G that are typically less than about 10 J m–2. Because the elastic properties of the polymer-coated elastomer and the polymer-coated indenter are different, one might expect that a mode II component (in-plane shear) of the stress field would drive the crack toward the low-modulus elastomeric layer [1]. While a mode II component of the stress field in the crack tip region is generally generated at the interface between materials with different mechanical properties, this is not true for the special case of a rigid indenter that is brought into contact with a thick, incompressible layer, as is the case with the elastomer substrate in our experiments [15]. In these tests, we assume that the contact stiffness of the elastomeric substrate is not substantially affected by the addition of the thin glassy or semicrystalline polymer layers. These layers deform by bending, and the load needed to bend the layers into conformity with the elastomer surface can be estimated from a simple plate deformation calculation. The deflection of the film can be modeled by the central loading of a disk [16], and comparing the stiffness of the disk to the contact stiffness of the elastomer substrate. For a circular plate that is loaded at the center and supported (but not clamped) at a radius of Rm, the stiffness is dependent on the modulus of the film, Ef, the thickness of the film, h, and its Poisson ratio m [Eq. (4)]. P 16pEf h3 d 3 mR2m 121

m

Ef h3 a2

4

The effective value of Rm must scale with the contact radius a, for an elastic half space, since this is the only length scale in the problem. An approximate equivalence between Rm and a gives the correct scaling as in Eq. (4). This stiffness can be compared, with the contact stiffness of the elastomer itself [see Eq. (3)], as in Eq. (5) where Es is Young’s modulus for the elastomer. P 8Es a Es a d 3

5

The relative contribution of the layer to the overall stiffness of the system is given by the ratio of the stiffnesses given by Eqs. (4, 5), i.e., (Ef/Es)(h/a)3. In our case this ratio is very small (Ef/Es & 1000 and h/a < 0.01), so that overlayers can be ignored in determining the compliance used to calculate the energy release rate. In the following sections, we describe the ability to run tests sensitive to values of G that are characteristic of processes confined to the interfacial region. First, the sample preparation and experimental procedure are described for two types of tests: a straight pull-off test and a cyclic interfacial fatigue test. Then, the results of these tests are presented with data for the forces required for fracture of the samples, as well as the calculations for G related to the cyclic fatigue test. We conclude with a discussion of the significance of the results that have been obtained for our model system.


23.2 Materials and Methods 23.2.1 Materials and Sample Preparation

As indicated in Fig. 23.1, a sample consists of a rigid glass indenter and an elastomer substrate of crosslinked poly(dimethyl siloxane) (PDMS), which are both coated with the polymer layers of interest. These layers include a semicrystalline layer of poly(ethylene oxide) (PEO) sandwiched between glassy polymer layers of poly(tetramethyl bisphenol A polycarbonate) (TMPC). These polymers will be described in more detail within this section along with the steps that were taken to select these polymers for the study of a glassy/semicrystalline interface. The poly(ethylene oxide) studied in these experiments was synthesized by anionic polymerization to give a polymer with a molecular weight of 70 kg mol–1 and a polydispersity of 1.07. PEO is a semicrystalline polymer that serves as a model system for the formation of crystalline structures at an interface because it has a low melting temperature of 66 8C. This polymer also has a Tg of approximately –60 8C in the amorphous phase [17, 18]. The low melting temperature of the PEO allows thin layers of the polymer to be easily spin-cast from warm solutions of butanol. In order to perform the specified adhesion tests, the PEO layer must be heated above its melting temperature while in contact with a glassy polymer layer to allow for mixing of polymer segments across the interface. Polystyrene (PS) is a glassy polymer that is well suited for these experiments because it can be chemically grafted to both the glass and elastomeric substrates. However, the relatively low glass transition temperature of PS (Tg = 100 8C) gives a narrow temperature window, between the PEO melting point and the glass transition temperature of PS, where suitable annealing treatments can be carried out. Also, thin layers of PEO dewet from PS upon heating and present an additional experimental difficulty. For this reason, we use poly (tetramethyl bisphenol A polycarbonate) (TMPC) as the glassy polymer layer in these tests. It is miscible with PS [19, 20], and has a glass transition temperature of 190 8C [19], which is well above the annealing temperature used in our experiments. The high glass transition temperature of TMPC ensures that the adhesion tests probe interactions that are due to the presence of the PEO.

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Fig. 23.2 Optical micrographs of thin PEO films (thickness = 80 nm) on TMPC substrates after annealing for 15 h at 85 8C. Dewetting is clearly observed on the PS substrates (a), but not on the TMPC substrates (b).

Dewetting studies were performed between PEO and TMPC as well as between PEO and PS to ensure that TMPC was an appropriate glassy polymer for studies of interfacial interactions. Fig. 23.2 shows optical micrographs from these dewetting tests. Two samples were made by spin-coating a layer, 130 nm thick, of TMPC or PS onto silicon wafers. An 80 nm layer of PEO was then spin-cast on top of each of these glassy polymer layers from solutions in n-butanol heated to approximately 60 8C. The samples were annealed in air for fifteen hours at 85 8C and then cooled to room temperature. The images were taken using a Spot RT Digital Camera attached to a Nikon Epiphot microscope. From these images, the dewetting of PEO from PS is clearly observed; whereas, no dewetting of PEO is observed when TMPC is used as the glassy layer. Differential scanning calorimetry (DSC) studies were also performed on TMPC/PEO polymer blends using a Mettler Toledo DSC 822 under a dry nitrogen atmosphere. Concentrated solutions with the following TMPC/PEO weight ratios were made in toluene: 90 : 10 TMPC/PEO, 50 : 50 TMPC/PEO, and 0 : 100 TMPC/PEO. The aluminum DSC pans were placed on a heated surface and the solution was dropped into the pan to allow evaporation of the toluene without crystallization of the PEO. Once the toluene had evaporated, the pan was removed from the heat and sealed. During the scan, the samples were initially heated from room temperature to 215 8C and held for 10 min to remove as much toluene as possible. A scan rate of 10 K min–1 was used as the sample was cooled from 215 8C down to –85 8C, held for 5 min, and then heated back to 215 8C. The 0 : 100 TMPC/PEO sample showed a strong endothermic melting peak at 66 8C. For the 50 : 50 TMPC/PEO sample this peak was shifted to 61 8C, and in the 90 : 10 TMPC/PEO sample this peak no longer existed. The trend seen in these tests indicates that as the percentage of TMPC in the blend is increased, the melting temperature associated with the PEO decreases. These re-


sults lead to the conclusion that PEO and TMPC are miscible polymers, which is discussed in Section 23.4.1. To make samples for the adhesion experiments, a crosslinked PDMS elastomer substrate was oxidized by a 30 min exposure to UV/ozone [21–24], using a Jelight Company Model 42 UVO-Cleaner. A thin layer of polystyrene (Mw = 38 000 g mol–1) with a trimethoxysilyl end group was spin-cast onto this oxidized elastomer from a solution in toluene. These functionalized chains were grafted onto the oxidized PDMS surface by annealing for 1.5 h at 125 8C, and the excess polystyrene chains were removed by rinsing the sample in toluene [25]. A layer of TMPC was then floated onto the rinsed elastomer surface and annealed for 1.5 h at 200 8C, which effectively grafted the TMPC to the surface because of the miscibility of TMPC and PS. This same process was repeated for a hemispherical glass indenter. The PEO layer was then applied to the TMPC layer on the elastomer substrate by spin-coating the polymer from butanol directly onto the TMPC-coated PDMS substrate. Because butanol is a nonsolvent for TMPC, contact between the butanol and the grafted TMPC layer did not result in any damage to the TMPC layer. By varying the concentration of the PEO solution, different thicknesses of PEO between 8 and 110 nm were deposited for adhesion testing. The corresponding thickness values for each concentration were determined by taking independent measurements of PEO layers spin-cast onto silicon wafers using a Tencor P-10 Profilometer. 23.2.2 Pull-Off Test

In order to perform the adhesion tests, the prepared samples were mounted in position in the heating chamber as shown in Fig. 23.3. The elastomer side of the sample was previously attached to a glass slide using a silicone glue and the slide was fastened into place on the sample holder, which was then connected

Fig. 23.3 Experimental set-up for the adhesion measurements.

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to a 50 g load cell in series with a Burleigh inchworm stepping motor. The load cell and stepping motor interfaced with a National Instruments LabVIEW program that allowed specification of the advancing and retracting velocity, as well as the maximum compressive load. The indenter was mounted on the underside of the lid of the heating chamber and aligned with the viewing window in the lid. This window allowed images to be captured during the test from the reflected-light microscope and camera mounted above the chamber. These images were used to measure the radius of contact at specific times during the test. The stepping motor advanced the elastomer sample into and out of contact with the indenter and a fiber optic displacement sensor was placed below the sample holder to detect the displacement associated with the movement of the sample. The load and displacement data that were obtained during the test were compiled with the LabVIEW program. Because the adhesion tests required that the surfaces be heated once they were brought into contact, the samples were held within a thermally insulated copper heating chamber. Dry nitrogen gas was passed through two in-line gas heaters, one on each side of the chamber, to heat the samples. A thermocouple was placed in the heating chamber in close proximity to the sample and was connected to a temperature controller so that the samples could be held at specific temperatures while in contact. The test was designed to be stress controlled during the heating and cooling of the samples, so that the samples were held at a constant compressive load while the displacement changed to account for thermal expansion of the device during the heating and cooling processes. The specific experiments that were performed using this method involved the interfacial interactions between PEO and TMPC. The TMPC/PEO-coated elastomer substrate was brought into contact with the TMPC-coated glass indenter. The samples were heated to 80 8C, which is above the melting temperature of the PEO. After 30 min of heating, the samples were cooled back to room temperature, and an increasing tensile displacement was applied until pull-off occurred. Typical data associated with such a test are plotted in Fig. 23.4. The load curve (Fig. 23.4 a) shows the load beginning at 0 mN and decreasing to –25 mN as the samples are brought into contact with one another. The load is fixed at –25 mN during the heating and cooling processes. The load then increases during the pull-off portion of the test and returns to zero after final separation of the two surfaces. The displacement curve (Fig. 23.4 b) illustrates the displacements recorded during this process. The large variation in displacement in the middle of the graph represents movement by the stepping motor to maintain a constant load and to compensate for thermal expansion during heating and cooling. The temperature curve (Fig. 23.4 c) shows a rapid increase in the temperature to 80 8C, followed by a plateau at this temperature and a slow decrease to room temperature. These tests were performed on TMPC layers with thicknesses of PEO ranging from zero (no PEO added) to 100 nm. Tests were also performed with a pre-annealing step before the sample and indenter were brought into contact. The sample and indenter were held out of contact at 80 8C for a specified period of time and then brought into contact to

23.2 Materials and Methods Fig. 23.4 Time dependence of (a) load, (b) displacement, and (c) temperature for a typical pull-off adhesion test. These graphs are compiled from pull-off tests for 130 nm TMPC layers with no PEO layer present.

–25 mN and held for 30 min. The samples were cooled back to room temperature and an increasing tensile displacement was applied until pull-off occurred. These tests were performed on TMPC layers with an 8 nm thick PEO layer applied to the elastomer side of the sample.

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23.2.3 Cyclic Interfacial Fatigue Test

The cyclic test was developed to limit the nominal stress experienced by the sample during testing. This test consisted of two parts. The first part included the initial heating and cooling cycles, and followed the same procedure as the pull-off test. The second part consisted of a cyclic tensile load being applied to the sample until failure. Typical data are shown in Fig. 23.5. Note that the pulloff portion of the test has been replaced by a cyclic variation in the tensile load between zero and 50 mN. During each increase of the load to 50 mN, the contact radius decreased by a small amount, corresponding to the growth of an interfacial crack. Continuation of the cyclic loading resulted in propagation of the interfacial crack until the coated indenter completely separated from the coated silicone elastomer.

23.3 Results

A series of control tests were run to investigate the interfacial interactions of a TMPC/TMPC interface after the layers were heated in contact. The TMPC layers for this test had thicknesses of 130 nm. A representative tack curve for this test is

Fig. 23.5 Time dependence of the (a) load, (b) displacement, and (c) temperature for the cyclic interfacial fatigue test.

23.3 Results Fig. 23.6 (a) Tack curve for a pull-off test involving two TMPC layers in the absence of PEO. The layers were heated to 80 8C while in contact and cooled to room temperature before pull-off. (b) Image of the elastomer side of the sample, showing the contact area after completion of the experiment.

shown in Fig. 23.6. The load and displacement begin at 0 mN and 0 lm, respectively, and then the load decreases to –25 mN. At this point, the heating and cooling of the sample occurred. The data for the pull-off proceeds from this point with a maximum tensile load, Pt, of 4.4 mN required for complete separation of the interface. This load is quite low, indicating that minimal adhesion occurred between the TMPC layers. The same tests were run with TMPC layers (all 130 nm thick) heated to different temperatures between 80 8C and 180 8C. Fig. 23.7 shows the maximum tensile pull-off loads associated with these tests. These data indicate that the adhesion between the layers remains minimal at least up to 150 8C. These

Fig. 23.7 Maximum tensile load versus temperature for separation of the TMPC/ TMPC interface with 130 nm TMPC layers on each side of the interface.

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low pull-off loads provide a control with which to compare the adhesion observed for TMPC/PEO/TMPC trilayer samples. Annealing temperatures of 80 8C were used for these subsequent experiments. PEO layer thicknesses for the trilayer samples varied between 8 and 110 nm. Data obtained for a sample with a PEO thickness of 12 nm are plotted in Fig. 23.8. The maximum tensile load in this case was 107 mN, and is indicative of the adhesive bonding attributable to the PEO layer. The adhesive strength of the interface was great enough to result in cohesive rupture of the PDMS elastomer, as illustrated by the optical micrograph of the contact area taken after the conclusion of the test. Similar tests were run for the other PEO thicknesses, and the maximum tensile loads for these different tests are shown in Fig. 23.9. The solid circle on the graph corresponds to a test that was run with a PEO layer that was not heated in contact. The pull-off load for this test was below 5 mN, indicating that the adhesion seen in other tests is a function of the samples being heated in contact. The other data points all have pull-off loads that are above 50 mN, demonstrating that the PEO layers produce an adhesive interface. However, most of these tests resulted in failure by rupture of the PDMS, instead of by interfacial crack propagation. Because of the irreproducibility of the details of this failure process, there is no noticeable trend in the pull-off loads associated with increasing PEO thickness. However, these pull-off tests are significant because they provide a basis for all the other tests run with this system. The control tests indicate that adhesion does not occur without PEO or with room temperature PEO, but that the presence of PEO heated in contact

Fig. 23.8 (a) Tack curve for a pull-off test similar to the one corresponding to Fig. 23.7, but with a 12 nm PEO layer between the TMPC layers. (b) Image of the elastomer side of the sample, showing the contact area after completion of the experiment.

23.3 Results Fig. 23.9 Maximum tensile loads as a function of the PEO layer thickness, for a series of pull-off tests where the samples had been heated to 80 8C while in contact.

with the TMPC creates an adhesive interface. These tests are also important because they indicate the limitations of the test by a statistical measurement of G. Although the results from the pull-off tests indicate that the interfaces in the trilayer samples were relatively strong, the maximum tensile loads were often controlled by the strength of the underlying PDMS elastomer, and not by the strength of the PEO/TMPC interfaces or by the cohesive strength of the PEO. Cyclic interfacial fatigue tests were used to explore the interfacial interactions in more depth by operating at loads that were below the critical load required for rupture of the PDMS. These tests also eliminated artifacts associated with possible meniscus formation of the PEO at the edge of contact between the TMPC layers. In the course of a cyclic fatigue test, edge effects are reduced by the propagation of an initial fatigue crack during the first cycle of the test. All subsequent cycles probe crack propagation in the region where the PEO was constrained between the glassy TMPC layers during the annealing treatment. Results for a cyclic test with a PEO thickness of 8 nm are shown in Fig. 23.10. A constant load amplitude of 50 mN was obtained for 40 cycles of the test. During the 40th cycle, complete separation of the interfaces occurred. The image included in Fig. 23.10 shows the interface after complete separation of the coated indenter and substrate. This image indicates that the failure occurred at the PEO/TMPC interface for most of the test, but that the final pull-off removed a layer from the center of contact. Because the missing layer can be observed visually, we can assume that both the PEO and TMPC were removed from the elastomer substrate. While there was layer removal in the final stages of the test, the interfacial separation that occurs during most of the cycles still provides information on the interfacial interactions between the PEO and TMPC. During this test, the displacement amplitude increased with each cycle. This increase in displacement is associated with crack growth and a decreasing contact radius, giving an enhanced compliance of the sample in accord with Eq. (3). As the crack propagated, the displacement increased with each cycle to maintain the same load amplitude.

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23 Contact Mechanics and Interfacial Fatigue Studies Between Semicrystalline and Polymer Films Fig. 23.10 Maximum and minimum loads (a) and displacements (b) for each cycle of the interfacial fatigue experiment. (c) Image of the contact area at the conclusion of the test.

The load and displacement data can be used in conjunction with the effective modulus of the system to determine the evolution of the contact radius during the test. This method is based on the use of the version of Eq. (3) given in Eq. (6). a

2DP 3KDd

6

In this equation, Dd and DP are the displacement and load amplitudes illustrated in Fig. 23.10. Use of Eq. (6) to calculate the contact radius can be automated, and reduces error that is likely to occur with visual measurement of the

23.3 Results

contact radii from the images. In Fig. 23.11 a, values of the measured contact radius are compared with values calculated according to Eq. (6). The excellent agreement between the measured and calculated values confirms the validity of this automated technique for determining the contact radius. Equation (1) can be used to calculate the applied energy release rate, G, for the maximum and minimum loads at each cycle. These values are plotted in Fig. 23.11 b. The final step in the analysis is to obtain the incremental decrease in contact radius between cycles, and to plot this as a function of DG, the difference in applied energy release rates between the maximum and minimum load conditions. The results of this analysis are shown in Fig. 23.12. The circular data points are associated with samples that were held at 80 8C for 5 min before cooling back to room temperature; square data points were for samples held for 10 min, and triangular data points for those held for 15 min. The lowest values of DG are obtained at the beginning of the experiment, and are determined by the load amplitude and by the contact radius that develops during the compressive portion of the experiment. For all three tests, a critical G occurs at about 3.5 J m–2. At this point, the incremental decrease in the contact radius during each cycle begins to increase substantially, leading to complete failure of the interface. The image included in Fig. 23.12 is an example of the elastomer surface after the test. The failure resulted in the removal of a layer that can be observed visually, much like the final removal of the layer seen in Fig. 23.10. Therefore, it

Calculated

Fig. 23.11 (a) Contact radius for each cycle of the fatigue experiment. (b) Maximum and minimum values of the applied energy release rate for each cycle of the fatigue experiment.

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23 Contact Mechanics and Interfacial Fatigue Studies Between Semicrystalline and Polymer Films Fig. 23.12 (a) Incremental decrease in the contact radius as a function of the difference in the maximum and minimum values of the energy release rate for each cycle of the fatigue experiment. Data are shown for tests with three different 80 8C contact times. (b) Image of the contact area at the conclusion of the 5 min test.

is believed that the PEO and underlying TMPC layers were removed from the elastomer. Because all three of these tests resulted in removal of these layers and because no trend is seen for these three time periods, it is believed that the diffusion of the PEO into the TMPC layer occurs at 80 8C before 5 min have elapsed. In fact, when samples are held in contact for even longer periods of time (greater than 25 min), failure again occurs by rupture of the PDMS, even with the cyclic test. These results suggest that the diffusion of the PEO into the TMPC occurs quickly and by 20 min of contact time it creates such strong interfaces between the layers that the PDMS failure strength is again the limiting factor. An additional set of pull-off tests was performed with a pre-annealing step incorporated into the test. During these tests, the samples were heated to 80 8C out of contact (pre-annealing) for a given period of time, then brought into contact at 80 8C for an additional 30 min. Once they had been cooled back to room temperature, the pull-off tests were performed. The values for G associated with these tests are shown in the graph in Fig. 23.13. These tests included pre-annealing times of between 3 and 16 min and the results indicate that there is a critical G of less than 1.5 J m–2 for annealing times greater than 5 min. These data also indicate that diffusion of PEO occurs quickly, as the adhesion drops between 5 and 10 min of pre-annealing time. The image included with this figure shows the elastomeric interface after the test. In these tests, no layer was removed from the interface. It is believed that the interfacial separation occurred between the top TMPC layer and the PEO layer because the diffusion of the PEO into the top layer of TMPC did not form a strong interface.

23.4 Discussion Fig. 23.13 (a) Energy release rates calculated for pull-off of pre-annealed samples heated to 80 8C for different pre-annealing times. After pre-annealing, samples were brought into contact at 80 8C for 30 min before cooling to room temperature and performing pulloff tests. (b) Image of the contact area at the conclusion of the 5-min test.

23.4 Discussion 23.4.1 Wetting Behavior and PEO/TMPC Miscibility

The ability of PEO to wet TMPC is an important result, given the widespread use of PEO surface coatings as biocompatible surfaces. Because PEO-modified surfaces have been found to reduce protein absorption and cell adhesion [26– 28], research has been conducted on the ability to attach PEO to materials used in implants [29–31]. Wettability of TMPC implies that the surface energy of PEO must be less than the surface energy of TMPC by an amount that is at least as large as the PEO/TMPC interfacial tension. Although the surface energy of TMPC is not directly known, studies by Kim et al. have shown that the surface energy of TMPC is greater than the surface energy of PS [32]. Studies by Sauer and Dee indicate that the surface energy of PS at 150 8C is 31.3 mN m–1 and that of PEO at 150 8C is 34.7 mN m–1 [33]. The surface energy for bisphenol A polycarbonate has been estimated as 35.1 mN m–1 at 150 8C [34]. Methyl groups are generally associated with low values of the surface energy, so we expect that TMPC will have a surface energy that is somewhat lower than the surface energy of bisphenol A polycarbonate. The ability of PEO to wet TMPC, and the presumed similarity in the surface tensions for these two polymers, indicate that PEO and TMPC are either thermodynamically miscible or very nearly so. Further evidence for the miscibility

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of these polymers comes from the character of the melting transitions observed in the TMPC/PEO blends. The behavior of these blends is similar to the miscible blend system of polyacetoxystyrene/PEO [18]. In both cases the melting temperature of the PEO component decreases with decreasing PEO content, and disappears altogether when the PEO content decreases below a critical amount. The decrease in the melting temperature that is observed can be explained by the decrease in the volume fraction of semicrystalline polymer in the system [35]. This behavior is consistent with the formation of a miscible blend, which for high TMPC content has a glass transition that is high enough for PEO crystallization to be kinetically hindered. Miscibility in the polyacetoxystyrene/PEO system is attributed to weak hydrogen bonding between the carbonyl groups on the polyacetoxystyrene and the methylene groups on the PEO [18]. A similar effect may be responsible for the miscibility in the TMPC/PEO system. PEO is in fact miscible over a large range of temperatures and compositions with bisphenol A polycarbonate, a polymer with a very similar molecular structure to TMPC [35]. 23.4.2 PEO/TMPC Interfacial Width and Adhesion

The development of adhesion between these two polymers can be related to the interfacial width between the polymers, as described in qualitative terms by Schnell, Stamm, and Creton [36, 37]. Very small values of the interfacial width are sufficient to give substantial adhesion; Gc values of &20 J m–2 were obtained for interfacial widths of only 4 nm. In our system, interfacial width, and hence the measured adhesion, are kinetically controlled, and are expected to depend on the temperature and contact time. Nevertheless, some insights can be gained by considering the values of the interfacial width that are obtained for immiscible polymers. For immiscible amorphous polymer pairs with a high molecular weight, the equilibrium interfacial width, w, is given by Eq. (7), where v is the Flory interaction parameter describing the segmental interactions and b is the statistical segment length [38]. 2b w p 6v

7

A corresponding expression [Eq. (8)], where v0 is the segmental volume, gives the interfacial tension, cab. cab

b kB T v0

r v 6

8

A useful quantity is the packing length, p, defined as the ratio of the volume of a polymer chain to its root mean square end-to-end distance, b [Eq. (9)] [39].

23.4 Discussion

p

v0 b2

9

Equations (1–9) can be combined to give Eq. (10) for the interfacial width in terms of the interfacial free energy and the packing length. w

kB T 3p cab

10

For PEO, p &1.9 Å [39]. Detailed information on the chain dimensions and density of TMPC is not available, but an estimate for p can be obtained from data for bisphenol A polycarbonate, which has p = 1.55 Å [39]. If the addition of four methyl groups is assumed to increase v0 by an amount that is proportional to the increase in molecular weight, while not affecting the statistical segment length, one obtains p = 1.9 Å for TMPC, a value equal to that of PEO. Equation (10) provides a link between the equilibrium interfacial width and the interfacial free energy for polymers with packing lengths similar to those of PEO and TMPC. If the polymers are fully miscible, as appears to be the case with the TMPC/PEO system, then cab = 0 and w = ?. In this case the actual interfacial width is kinetically limited, and is determined by the contact temperature and contact time. In our experiments the contact temperature lies between the glass transition temperature of TMPC (190 8C) [19] and PEO (–60 8C) [17, 18]. The dynamics of interfacial broadening in this case are controlled by the plasticization of the TMPC by the PEO. Large values of the interfacial toughness in glassy polymer systems are obtained when the interfacial width becomes comparable with the entanglement spacing. In the work of Schnell, Stamm, and Creton the smallest interfacial width of 4 nm was about half the entanglement spacing, yet the measured interfacial toughness for this sample was already 20 J m–2 [37]. Kinetic control of the interfacial broadening enables us to access very low values of the interfacial width and toughness that are difficult to obtain in an equilibrium experiment. Two separate groups of experiments have been performed to explore the kinetics of the TMPC/PEO interaction. In the first set of experiments, the samples were held in contact at 80 8C for different lengths of time at a specific compressive load. In the second set, the samples were pre-annealed out of contact at 80 8C for different lengths of time before they were brought into contact at a specific compressive load. Fig. 23.14 shows a schematic of the proposed interaction of the PEO chains with TMPC in these two types of tests. In the first set of tests, the adhesion that develops is a result of the diffusion of the PEO into both sides of the TMPC, resulting in plasticization of the TMPC, or an effective reduction in the Tg of these TMPC chains. Plasticization provides enough mobility for the chains to reorganize and form a strong interface. In the second set, the diffusion of the PEO into one side of the TMPC leaves the top layer of TMPC unaltered, with a high Tg. When the samples are then brought into contact, the PEO chains do not have the same ability to penetrate the TMPC as

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23 Contact Mechanics and Interfacial Fatigue Studies Between Semicrystalline and Polymer Films Fig. 23.14 Schematic of PEO chain interaction with TMPC based on heating conditions. Heating in contact results in an adhesive interface, whereas pre-annealing reduces the adhesion at the interface.

they did when heated in contact. While the pre-annealing step allows plasticization to occur between the PEO and the lower TMPC layer, the glass transition at the surface of these layers has increased from the initial Tg associated with PEO. Therefore, when the top TMPC layer is brought into contact with the already plasticized PEO/TMPC layers, the mobility required for the formation of an adhesive interface at the upper TMPC contact no longer exists. The net result is a reduction of adhesion as the pre-annealing time increases. 23.4.3 PDMS Rupture

Failure of the bulk PDMS substrate was observed in many of the axisymmetric adhesion tests. In order to promote failure through crack propagation at the interface, the test must be designed so that the applied energy release rate remains below the critical value corresponding to cohesive rupture of the silicone elastomer. This limitation introduces the importance of an interfacial fatigue test, where crack propagation occurs under cyclic loading conditions at relatively low values of G. Our tests were performed with a constant load amplitude, so that G increases as the contact radius decreases, as given by Eq. (1). With our current approach, G eventually increases to a value that is large enough for cohesive rupture to occur. The energy release rate can be held constant or adjusted in other ways by using the displacement signal to determine the sample compliance and contact radius, and adjusting the load during the test to maintain the desired energy release rate. Future studies will incorporate these more complex loading histories to study the effects of temperature, contact time, and PEO layer thickness on the interfacial adhesion.

References

23.5 Conclusion

The studies of PEO and TMPC interactions provide insight into the compatibility of PEO on TMPC and the adhesive forces that develop at the interface between these two polymers. PEO layers with a thickness of 12 nm provide an interface of sufficient strength to rupture the underlying silicone elastomer. Adhesion develops only when the materials are heated above the PEO melting temperature while in contact with the TMPC, a result that is attributed to the timedependent mixing of the thermodynamically miscible combination of the rubbery PEO and the glassy TMPC. In order to avoid complications associated with the rupture of the silicone elastomer, interfacial fatigue tests were developed to study the fatigue resistance of the PEO/TMPC interface under cyclic loading conditions. The contact radius is monitored in these experiments by simultaneously measuring the load and displacement amplitudes. In this manner a single experiment can be conducted to obtain the crack advance per cycle as a function of a range of loading conditions.

Acknowledgments

This work was supported by the MRSEC program of the National Science Foundation (DMR-0076097) at the Materials Research Center of Northwestern University. We also thank Mark Hammersky of Proctor and Gamble for providing the TMPC used in these experiments. Much of this work was published previously in Ref. [40].

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24 Local and Global Aspects of Adhesion Phenomena in Soft Polymers M.-F. Vallat

Abstract

In assemblies of soft polymers such as elastomers, interdiffusion and co-crosslinking are competitive mechanisms. They lead to the properties of the interphase at the molecular level, which is responsible for the stress transfer. However when elastomers are considered, crosslinking is generally obtained by the addition of crosslinking agents. It is therefore necessary to take into account the migration of the crosslinking agent, which leads to a gradient of properties.

24.1 Introduction

When two polymers are brought into contact under conditions of sufficient mobility and compatibility, the interdiffusion of the chains is responsible for the disappearance of the original interface. For elastomers, this rule applies even at room temperature, which is above their glass transition temperature. The mobility of the chains of elastomers being high, both at and above room temperature, chain extraction is very often quite easy unless separation is performed at a high rate or a low temperature, in which case failure of nonextracted chains can occur. This means that chemical crosslinks most probably have to be formed in this molecular interphase to obtain strong joints. Because generally elastomers are crosslinked during joint formation, interdiffusion of the chains at the interface and formation of crosslinks in the molecular interphase are competitive mechanisms [1]. Both mechanisms are necessary to ensure high interfacial strength. However, chain mobility is rapidly restricted by the formation of the three-dimensional network. Therefore, the formation of elastomer joints is generally a complex phenomenon because interfacial development and network formation occur simultaneously. The properties of the interphase at the molecular level is responsible for the stress transfer. Its spatial extent and the small number of crosslinks created in this area make it very difficult to obtain a local description of the properties. Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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24 Local and Global Aspects of Adhesion Phenomena in Soft Polymers

Moreover, when the elastomer networks are obtained with the help of crosslinking agents, an additional contribution has to be considered in the study of the assembly behavior. Indeed, the crosslinking agents can migrate easily across the interface and induce modifications of the mechanical properties of the material at the interface and up to some distance from it. The gradient of mechanical properties which is created is important for the locus of failure as well as the dissipated energy during measurements of the interfacial strength. The characterization of the elastomer network at a fine or mesoscopic scale, i.e., from the interface to a distance of about 500 lm away, is an important challenge. In this paper, we will review how the two aspects in elastomer joints have been considered in various work that has been done in our institute.

24.2 The Molecular Interphase

In this section, we will consider homogeneous, symmetric elastomer joints for which we want to find the relationship between peel strength and density of bonds in the molecular interphase. The following definitions are used: · homogeneous elastomer joints are joints in which the same elastomer formulation is present on both sides of the interface; · symmetric elastomer joints are joints which are obtained by joining two elastomer sheets with the same rubber formulation and the same degree of crosslinking (including uncrosslinked sheets). With such joints, it is theoretically possible to vary the density of interlinks between zero and that corresponding to the crosslinking density of the bulk material. Indeed, when the elastomer sheets are assembled in the uncrosslinked state (the chains are free), the mobility is high and the crosslinking agents are fully available to create crosslinks both in the bulk of the material and in the interfacial region during the crosslinking procedure. The interfacial strength should then be identical to the cohesive strength of the elastomer at the same degree of crosslinking. At the other extreme, when the network is formed before the assembly step, no crosslinking agent is left in the bulk material and no chemical bond can be created during the contact step. The interfacial strength is only obtained through van der Waals interactions and in some circumstances through interdiffusion of chain ends or pendant chains. If one uses the two-step assembly procedure proposed by Chang and Gent [2], intermediate states can be reached. For such a case, the two elastomer sheets are partially crosslinked separately (i.e., in the pre-crosslinking step). The sheets are then brought into contact and a subsequent crosslinking step (i.e., the postcrosslinking step) takes place. Conditions of pre- and post-crosslinking are chosen in such a way that the two steps always lead to the same final degree of


crosslinking, independently of the degree of pre-crosslinking. The density of elastic chains m reached after each step is evaluated by techniques such as swelling or modulus measurements. The number of crosslinks Dm formed during the assembly step is then given by Dm = mpost–mpre. It is generally accepted that the theoretical number of interfacial crosslinks is equal to Dm (bulk values). This procedure allows one to compare the peel energies quantitatively, because the bulk viscoelastic properties of the materials in the joints are identical. However several questions can be raised. Interdiffusion and co-crosslinking mechanisms are in competition during the joint formation so that the same density of crosslinking may not be reached in the interfacial zone as in the bulk. Therefore, the procedure we propose for separating these two mechanisms is electron beam crosslinking: In the first step of contact, chain diffusion can take place without changing the density of crosslinking; in the second step, the density of crosslinking is modified without changing that state of diffusion, by irradiation at room temperature for a short time. The question of the relationship between the interdiffusion depth and interfacial co-crosslinking density can be faced. We will now report the results of autohesion for homogeneous, symmetric joints of polyisoprene rubber (IR) and styrene–butadiene copolymer (SBR) both vulcanized by a sulfur-based system (Section 24.2.1), and of ethylene–propylene diene terpolymer (EPDM) crosslinked by an electron beam (Section 24.2.2). 24.2.1 Autohesion of Polyisoprene

Autohesion of polyisoprene rubber (Natsyn 2200, a synthetic high-cis-1,4-polyisoprene) and styrene–butadiene copolymer has been studied. Both elastomers were reinforced by carbon black and crosslinked by a sulfur-based system (see Table 24.1) [3]. The glass transition temperatures of the elastomers were not significantly changed by crosslinking and were equal to –66 8C and –53 8C for the IR and SBR, respectively, as measured by DSC analysis.

Table 24.1 Formulation of vulcanizates of IR and SBR [phr]. a)

Elastomer Other constituents Carbon black N 347 Stearic acid Antioxidant (6-PPD) Sulfur Accelerator (CBS) Zinc oxide a)

100 50 1.5 1.5 1 1 4

phr = parts per hundred of rubber by weight. CBS = N-cyclohexyl-2-benzothiazole-sulphenamide. 6-PPD = N-(1,3-dimethylbutyl)-N'-phenyl-p-phenylenediamine.

389

390


Two temperatures of crosslinking, 126 8C and 150 8C, were investigated. The state of crosslinking could be followed by swelling measurements at equilibrium. The elastomers as received dissolved in a good solvent. After the crosslinking reaction, the amount of solvent absorbed by the network depended on the molecular weight between crosslinks. This amount could be weighed. Lorenz and Parks [4] have shown that 1/Q2p (where Qp is the swelling ratio defined as the ratio of the mass of solvent absorbed at equilibrium to the mass of elastomer after solvent swelling and drying) varies as the density of crosslinking. Cyclohexane was used for both polymers. The values of 1/Q2p are plotted as a function of time of crosslinking for both the temperatures considered in Fig. 24.1, which shows the expected effects due to sulfur vulcanization when temperature of crosslinking is higher: · the induction time decreases (for shorter times no three-dimensional network is formed); · the kinetics of crosslinking is more rapid; · the degree of crosslinking at the optimum of crosslinking, as given by the maximum of the curves, is lower.

Table 24.2 Induction time [min] for the two elastomers and the two temperatures of crosslinking.

Temperature of crosslinking [8C]

IR SBR

126

150

15 37

3 9

Fig. 24.1 Degree of crosslinking as given by the swelling ratio as a function of time of crosslinking for the two elastomers and the two temperatures of crosslinking.


This effect is more pronounced for SBR because crosslinking by radicals can also occur for this elastomer. No reversion is observed over the times considered. The assemblies were prepared according to Gent’s procedure in two steps as described above. To obtain various states of crosslinking before assembly, the time of pre-crosslinking was changed. The degree of pre-crosslinking a was determined from Eq. (1). a

Qp 2 tpre Qp 2 topt

1

The total time of crosslinking was maintained constant and equal to the optimum time of crosslinking. At the optimum of crosslinking (tpre = topt), a was equal to 1: There was no more free crosslinking agent left and the number of crosslinks which could be created during the contact was nil. When there was no pre-crosslinking (tpre = 0), the density of crosslinks formed during contact was equal to the total value corresponding to the recipe given in Table 24.1. The two temperatures of crosslinking were used. If the elastomer was pre-crosslinked at 126 8C, the joint was crosslinked at the same temperature (the same was valid for 150 8C). A 1808 peel test was used to evaluate the performance of the adhesive joints. The results presented here were obtained at room temperature and a peel rate of 0.5 mm min–1; this rate corresponds to reduced energy dissipation. The specific peel energy is plotted as a function of the degree of pre-crosslinking a in Fig. 24.2. The effect of the crosslinking temperature was higher for SBR joints than for IR joints. The influence of the bulk properties is illustrated by the high peel strength of the SBR joints. This can be related to the loss modulus of SBR,

Fig. 24.2 Peel strength G versus degree of pre-crosslinking a for the two elastomer joints crosslinked at 126 and 150 8C, respectively.

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392


which was about twice that of IR at room temperature (measured at 250 Hz). Dissipated energy during the separation of the two sheets of elastomer was therefore higher [5]. It should also be mentioned that the ratios of the elastic and loss moduli of SBR (E'126 8C/E'150 8C and E''126 8C/E''150 8C) at room temperature were 1.2, whereas they were 1.1 and 1 respectively for IR. Below certain a values, peeling was no longer possible because of the limited reinforcement by the backing cloth and the strength of the backing cloth itself. These a values were higher for a crosslinking temperature of 150 8C for both elastomers, which means that higher interfacial elastomer–elastomer strength was reached even though the degree of pre-crosslinking was closer to the optimum. A comment should also be made on the locus of failure: for IR, it was apparently interfacial (both peeled parts appeared smooth with a glossy aspect, as before assembly) for peel strengths up to about 2.5 kJ m–2, and cohesive above that value (the separation led to a more or less rough surface and finally became impossible). For SBR joints, the locus of failure also changed and an additional type of failure was observed. Samples showed stick–slip behavior between 7 and 15 kJ m–2: The peel force oscillated between a minimum and a maximum value. During the induction time (Table 24.2), no crosslinks were formed, and the polymer remained soluble in a good solvent. This time was about four times as long for 126 8C as for 150 8C. On the one hand, the intrinsic mobility and the diffusion coefficient were both lower. However, the mobility of the chains remained high for a longer time before network formation occurred, which means that the chains were progressively integrated to the network and became immobile. The thickness e of the interdigitated region is given as a function of contact time t by Eq. (2), where R0 is the radius of gyration of the chain, trep is the time that the chain needs to leave its initial tube, L is the contour length of the chain, and D is the curvilinear diffusion coefficient [6]. et R0

t trep

1=4 and

trep

L2 2Dt

2

Diffusion coefficients are proportional to 1/M, the molecular weight of linear chains [7]. They are not well known and therefore neither is the interdigitated thickness. Hence, it is not possible to say whether the observed behavior has to be related to the interdiffusion depth, or to the number of crosslinks formed in the interfacial region or, most likely, to both effects. These results given for the elastomer joints crosslinked by a sulfur-based vulcanizing system show that it is very difficult to separate interdiffusion and crosslinking mechanisms because the temperature influences both the chain mobility and the kinetics of network formation. An original method for separating these two mechanisms has been used for studying the interfacial strength of EPDM joints [8]. It is based on appropriate conditions for the joint formation leading to independent control of the two mechanisms. Crosslinking by an electron beam (2.2 MeV energy, providing a homogeneous radical density in samples up to 2 mm thick) can be used to ob-


tain the polymer network at around room temperature without any addition of chemical species. It is then possible to control diffusion of the chains by controlling both the temperature and contact time, and to superimpose the macromolecular network without modifying the state of interdiffusion. 24.2.2 Autoadhesion of EPDM

Some preliminary results were obtained for IR and SBR joints [9], but EPDM autohesion was studied more extensively [8]. EPDM with 53 wt.% of ethylene, 42% of propylene, and 5% of ethylidene norbornene was used. Its glass transition temperature was –55 8C. Two sheets of this elastomer (each 1 mm thick) were crosslinked separately by dicumyl peroxide. The network characteristics were: m = 10 mol m–3; sol fraction = 25 wt.%. The crosslinking time was chosen so that no peroxide was left after this first step. The two sheets were then joined for various periods of time at room temperature. Chain interdiffusion was occurring during this second step without any change in the state of crosslinking. After being joined under various conditions, the joints were subjected to an electron beam for the third step of the joint preparation, corresponding to the crosslinking of the elastomers. In our experimental conditions (2.2 MeV accelerator – Model AS2000 from HVEC at Aerial, Strasbourg, France), the joint was submitted to a dose of 130 kGy at room temperature [10]. With a dose rate of 1.4 kGy s–1, the crosslinking step lasted only about 100 s and the temperature reached by the sample due to the reactions was below 50 8C. The degree of crosslinking m reached in the bulk after the electron beam treatment was 105 mol m–3 with a sol fraction of 8 wt.% [11]. It can therefore be considered that only crosslinking was occurring during this third and last step of joint formation, without any change in the depth of interdiffusion. The adhesive strength of post-crosslinked samples was evaluated by a 1808 peel test which measured the force of separation at a constant peel rate (5 mm min–1) at room temperature. The results reported in Fig. 24.3 show the effect of contact time (ranging from 1 h up to one month) in the second step of joint formation at room temperature. Up to approximately 300 h of contact, no effect of time is seen. For longer times, the peel strength increases slightly. Although the scatter is important, this weak trend can be related to crosslinks formed by irradiation in the molecular interphase. This is confirmed by the fact that no separation is observed for these joints immersed in a good solvent such as cyclohexane, whereas spontaneous delamination is observed for shorter contact times. Therefore, the presence of crosslinks in the interdiffused depth is clear after long contact times but no crosslinks, or only a few, are formed at shorter times. Because the density of crosslinks inside the elastomer sheets is the same in all cases (the same dose is received), only the effect of diffusion depth on this number of crosslinks in the interfacial domain can explain the observed behavior. Moreover, Eq. (2) tells us that the thickness of the interdigitated zone in-

393

394


Fig. 24.3 Peel strength versus contact time of contact at room temperature for EPDM joints made of partially crosslinked networks which are post-crosslinked by an electron beam (peel rate = 5 mm min–1).

creases continuously with time. Therefore it can be proposed that a critical thickness has to be reached to give efficient interfacial crosslinks. When the electron beam crosslinking was performed after the contact steps at temperatures higher than room temperature (ranging from RT up to 170 8C), the peel energy increased continuously with the temperature of the contact step (Fig. 24.4). It finally reached a value close to the tear strength of the bulk material at the same degree of crosslinking. The results indicate that the number of crosslinks which could be created in the interdiffused zone increases with the temperature of contact. About an hour or so at 80 8C is as efficient for the interfacial strength as about 300 h at room temperature. This confirms the very low diffusion rates of the chains at the interface and is in agreement with results reported by others [12, 13] and also by us [14]. The diffusion coefficients depend strongly on molecular weight and also on branching. They decrease drastically when branching is present, as could be shown by comparing the behavior of linear and branched polymer chains [15]. It is known that the addition of the diene terpolymer introduces such branches into EPDM [16] and we then expect low diffusion coefficients, even for the chains before they are included in the network. These results confirm that the number of interfacial crosslinks depends on the interdiffusion depth and that a critical depth has to be reached to obtain efficient crosslinks for high interfacial strength.

24.3 Macroscopic Interphases

Fig. 24.4 Peel strength versus temperature of contact for EPDM joints made of partially crosslinked networks which are post-crosslinked by an electron beam (peel rate = 5 mm min–1).


When nonsymmetric joints (with the same polymer but different crosslinking recipes or different crosslinking states) or nonhomogeneous joints (different polymers) are made with polymers containing crosslinking agents, an additional contribution to the adhesive behavior has to be considered. Indeed, migration of crosslinking agents or small molecules can occur very easily over long distances before or even during the crosslinking reaction. This effect can lead to macroscopic interphases with a gradient of mechanical properties which play an important role in the measured separation strength and the locus of failure observed during propagation of the failure in assemblies. Two examples will be used to describe this aspect. 24.3.1 Vulcanized Elastomers

Preliminary experiments were done on adhesion of the two elastomers considered in Section 24.2.1, polyisoprene and styrene–butadiene copolymer [9]. One sheet was fully crosslinked before contact. Then the joints were made with the second uncrosslinked sheet. No significant variation of the degree of crosslinking (and therefore no reversion) was observed when the elastomers were crosslinked for longer times in the conditions given in Table 24.3. In the nonsymmetric joints of IR, very high strengths were reached, but could not be measured because of the limited strength of the backing cloth. For the SBR nonsymmetric joints, failure occurred in the (SBR)2 sheet.

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396

24 Local and Global Aspects of Adhesion Phenomena in Soft Polymers Table 24.3 Conditions of assembly of various joints and the resulting peel strength.

Joint

SBR/SBR IR/IR

SBR/IR IR/SBR

Crosslinking conditions

Peel rate [mm min–1]

Peel strength [kJ m–2]

5.9 11.0 failure of the backing

tpre [min]

tpost [min]

(SBR)1 (SBR)2 (IR)1

150 0 110

150 150 110

0.5 5 0.5

(IR)2 SBR IR IR SBR

0 0 110 0 150

110 150 150 110 110

5 0.5 5 0.5 5

9.6 14.0 1.6 2.2

In the nonsymmetric, nonhomogeneous joints (SBR/IR and IR/SBR), failure occurred in the SBR sheet but the peel strength compared at the same peel rate varied significantly. The existence of a gradient of properties due to the migration of sulfur or sulfur intermediates was then proposed [9]. Nonsymmetric joints of polyisoprene (Fig. 24.5 and Table 24.4) were studied in order to gain a better understanding of this effect [17, 18]. Two sulfur-based vulcanizing recipes, a “conventional” one (C) and an “efficient” one (B), were kept as simple as possible: sulfur, accelerator, activator, and anti-oxidant. The amounts of sulfur and accelerator were adjusted to give about the same level of stiffness for both components and to be comparable with the semi-efficient system used in Section 24.2.1 (Table 24.1). These conditions would promote migration of the vulcanizing agents through the interface due to their high concentration gradient during joint formation. It is then possible to study its effect on the measured strength. When these two sheets were assembled, the interfacial strength could not be measured because the interface was very strong. The two mechanisms, interdiffusion of the chains and co-crosslinking, both occurred. A modulus gradient

Fig. 24.5 Nonsymmetric joints B and C of polyisoprene (r = [sulfur]/[accelerator]).


Fig. 24.6 Modulus as a function of distance from the interface without (tc = 0 h) and with preliminary contact (tc = 48 h) before co-crosslinking.

was observed through the thickness of the sample; it could be measured by a microindentation technique (Fig. 24.6). A Fischer instrument with a spherical indenter of 400 lm diameter was used to measure the penetration as a function of the applied force at constant rate. The compression modulus could then be determined [18]. There was an important increase in the modulus over a large distance from the interface, due to migration of free sulfur as well as of accelerator and their complexes which do not react with the IR chains [19]. Contact at 40 8C for 48 h before co-crosslinking changed the gradient of the modulus, as can also be seen in Fig. 24.6: r, the ratio of sulfur to accelerator, differed under the two sets of conditions and varied over the sample thickness, leading to the modulus gradient observed. The introduction of a gradient of mechanical properties could also be observed when partially crosslinked sheets containing free sulfur or accelerator were joined. For example, sheet B was pre-crosslinked to a = 0.9 and no free sulfur was left. Sheet C was pre-crosslinked up to a = 0.7 and a significant amount of free sulfur was left. The joint between these two sheets was then heated for an additional 68 min at 130 8C. Two samples have been compared: The two sheets were post-crosslinked immediately after joining, or after contact for 48 h. When the peel energy of these joints was measured, the sulfur migration had an effect on the measured value as expected – see Fig. 24.7, in which the peel energy is given as a function of peel rate. The interpretation of the experimental results in terms of interfacial co-crosslinks becomes very hazardous.

397

398


Fig. 24.7 Peel strength G versus peel rate R: influence of the contact time at 40 8C after pre-crosslinking of the sheets separately and before post-crosslinking.

24.3.2 Polyurethane Joints

The highly filled polyurethane joints that were studied are described in Fig. 24.8 [20]. The polymer chains on both sides of the interface are hydroxytelechelic polybutadiene chains (HTPBs). On the one hand, we have PUe in which the main additives are diisocyanate (isophorone diisocyanate) and 80% of reinforcing fillers (aluminum and potassium chloride). On the other hand, PUa contains the same diisocyanate and only 20% of reinforcing filler (carbon black),

Fig. 24.8 Nonsymmetrical assemblies of polyurethanes.


but short-chain diols are also added as chain extenders. A PUa network is constituted by hard and soft domains. The hard domains, which play the role of physical crosslinks, are due to hydrogen bonding between segments based on short-chain diols and diisocyanate [21]. The infrared spectra obtained by the ATR method after completion of the crosslinking reaction of these materials (Fig. 24.9) show the typical band of the hydrogen-bonded carbonyl groups around 1695 cm–1 and free carbonyl groups (1725 cm–1). The degree of crosslinking of PUa material can be changed by varying the time of crosslinking at 60 8C prior to contact with PUe and complete crosslinking of the joints (14 days). The interfacial strength is measured in a 90 8C peel test. The strength depends on the time of pre-crosslinking but, in all cases, failure always occurs in PUe at some distance from the original interface, which could be easily identified because of the colors of the fillers. The material after separation was characterized by ATR spectroscopy (Fig. 24.9). The carbonyl peak in PUe had a maximum around 1734 cm–1 (curve 5) in the bulk of the material. After the peel test, the carbonyl peak in PUe but close to the interface (curves 2–4) was no longer distinguishable from that in PUa (curve 1) at about 50 lm from the interface. Moreover, a contribution of hydrogen-bonded C=O groups appeared even at a distance of less than 50 lm. The locus of failure depends on parameters such as the state of crosslinking of PUa before contact and the characteristics of the HTPB used (molecular weight distribution, functionality, etc.) but also on testing conditions (peel rate, air or liquid). It can be related to the properties of the material close to the interface. The shape of the carbonyl peak can only be explained by migration of short segments (short-chain diols or short segments constituted by diols and diisocyanate) from PUa toward PUe during joint formation. Layers of material 50 lm thick were cut parallel to the interface by using a microtome at –60 8C. Tensile properties were measured at a rate of 1 mm min–1 at room temperature and the

Fig. 24.9 Evolution of the C=O peak at various distances (shown in the Fig.) from the interface after the peel test.

399

400


Fig. 24.10 Evolution of the modulus of layers 50 lm thick, cut at various distances from the interface, as a function of the distance from the interface.

modulus was determined from the first part of the curve. The variations in modulus close to the original interface over about 200 lm on both sides are shown in Fig. 24.10. They confirm the migration of short segments which are able to be hydrogen-bonded and create hard domains or physical crosslinks in PUe and increase in modulus. This modification comes with a diminution of E in PUa.

24.4 Conclusion

The competition and the relationship between the two main mechanisms, interdiffusion of the chains across the interface and co-crosslinking in the interdiffusion domain, have been shown by means of two examples. The interpenetration length determines the efficiency of the crosslinking reaction. Although the strength of this interface or interphase at the molecular level is of prime importance in the adhesive behavior of elastomer joints, another aspect needs to be considered due to the migration of the crosslinking agents over large distances leading to a gradient of mechanical properties. This gradient contributes to the measured strength but also to the locus of failure.

References

References 1 A. Aradian, E. Raphaël, P. G. de Gennes, 2 3 4 5 6

7 8

9 10

11

Macromolecules, 2000, 33, 9444 R. J. Chang, A. N. Gent, J. Polym. Sci., Phys. Ed., 1981, 19, 1619 M. F. Vallat, M. Stachnik, J. Schultz, J. Adhesion, 1996, 58, 183 O. Lorenz, C. R. Parks, J. Polym. Sci., 1961, 50, 299 J. Schultz, A. N. Gent, J. Chim. Phys., 1973, 70(5), 708 F. Brochard-Wyart, in “Fundamentals of Adhesion”, Lieng-Huan Lee Ed., pp 181– 205 (1991) P. G. de Gennes, J. Chem. Phys., 1971, 55, 572 F. Ruch, M. O. David, M. F. Vallat, J. Polym. Sci.: Part B: Polym. Phys., 2000, 38(23), 3189 M. Stachnik, PhD University of HauteAlsace (91-Mulh-0173) 1991 A. Strasser, F. Kuntz, E. Marchioni, R. Seltz, Nucl. Instr. and Meth. B, 1991, 56/57 1223 M. F. Vallat, F. Ruch, M. O. David, Eur. Polym. J., 2004, 40(7), 1575

12 C. M. Roland and G. G. A. Böhm, Macro-

molecules, 1985, 18, 1310 13 T. Q. Nguyen, H. H. Kausch, K. Jud, M.

Dettenmaier, Polymer, 1982, 23, 1305 14 M. O. David, T. Russ, M. F. Vallat, R.

Brenn, in preparation 15 C. B. Gell, W. W. Graessley, L. J. Fetters,

16 17 18 19 20

21

J. Polym. Sci.: Part B: Polym. Phys., 1997, 35, 1933 B. J. R. Scholtens, Rubber. Chem. Technol., 1984, 57, 703 S. Giami, PhD University of Haute-Alsace, Mulhouse – France, 1998 M. F. Vallat, S. Giami, A. Coupard, Rubber Chem. Technol., 1999, 72, 701 D. S. Campbell, Rubber Chem. Technol., 1971, 44, 771 M. F. Vallat, N. Bessaha, J. Schultz, C. Combette, J. Maucourt, J. Appl. Polym. Sci., 2000, 76(6), 665 Z. S. Petrovic, J. Ferguson, Prog. Polym. Sci., 1991, 16(5), 696

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25 Calibration and Evaluation of Nonlinear Ultrasonic Transmission Measurements of Thin-Bonded Interfaces S. Hirsekorn, A. Koka, S. Kurzenhäuser, and W. Arnold

Abstract

Interfaces in bonded structures influence the mechanical behavior of components significantly, and often limit their load capacity. This requires nondestructive testing techniques allowing one to investigate the interaction forces in adhesive joints and to evaluate the quality of bonds. To this end the nonlinear stress–strain relationships of adhesives and adhesive interfaces, which cause a nonlinear modulation of ultrasonic waves in reflection as well as in transmission, may be exploited. Bonded interfaces which are much thinner than the ultrasonic wavelength can be approximately described only by the binding forces, without explicitly taking into account the material properties of the adhesive layers. These may be measured by the amplitudes and phases of ultrasonic waves transmitted through the interface. Measurements are presented on aluminum plates joined together by thin epoxy adhesive layers. A threshold behavior of the harmonics generated in the adhesive layer has been observed. Their amplitudes depend on the excitation following the power series expansion of a quasi-static interaction force curve, and their phases vary little for low-amplitude excitation. Exceeding the threshold causes a change in the response of the interface. The input and output ultrasonic amplitudes in the interface are calibrated interferometrically to obtain the absolute interaction force. The ultrasonic transmission data are related to destructive tensile tests of the adhesive bonds.

25.1 Introduction

The interfaces in bonded structures influence the mechanical behavior of components significantly. Therefore, an important task in nondestructive testing (NDT) is the investigation of the interaction forces in adhesive joints and the development of techniques to evaluate the bond quality. The load capacity of such joints is often limited by regions of weak bonding. As in all materials, the Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

404

25 Calibration and Evaluation of Nonlinear Ultrasonic Transmission Measurements

stress–strain relationship of the adhesives and adhesive interfaces becomes more and more nonlinear with increasing strain [1]. Assuming that the ultrasonic strain is sufficiently high for the nonlinear part of the stress–strain curve to be covered, then a nonlinear modulation of ultrasonic waves will occur, resulting in the generation of higher harmonics (and possibly also subharmonics) in both reflection and transmission. The amplitudes and phases of these waves contain information about the stress–strain curve of the adhesive layer and hence the interaction forces transferring the ultrasound. There have been a number of previous attempts to relate the appearance of higher harmonics in the transmission of ultrasound through bonded structures to the quality of the bonds [2–9]. Commonly used is the so-called nonlinearity parameter b2, a measure of the generation of only the second harmonic [2], and the distortion factor K which describes the complete nonlinear content of the response [7]. In this paper calibrated measurements on samples consisting of two aluminum plates joined together by a thin epoxy layer are presented and discussed. The amplitudes and phases of the ultrasonic waves transmitted through the bond are considered. The measurements are related to the results of destructive tensile tests of the adhesive layer. For the calculation of the transmission and reflection coefficients for ultrasonic waves at thin-bonded interfaces it is sufficient to take only the binding forces into account, allowing one to calibrate the stress and strain levels acting on the interface [3, 7–10]. This in turn opens the way to obtaining the gross local interaction forces by measuring the amplitudes and phases of the transmitted ultrasonic waves [10]. Damping and hysteretic effects are included. Fracture mechanics, not considered in this context so far, must be taken into account in the presence of delaminations. We are aware of the fact that the tensile strength of adhesives depends on the load rate. If we assume that the load rate e_ for ultrasonic displacements is approximately the strain e multiplied by the ultrasonic frequency f, e_ = e f , we obtain for the example e = 10–4 and f = 2 MHz the load rate e_ 2 102 s–1. At such high load rates, the tensile strength of adhesives can be about two to five times higher than in quasi-static tensile tests with a strain load rate e_ of 10–2 s–1 [11].

25.2 Experimental and Calibration Procedure

The experimental set-up has been described previously [3, 7, 8]. Briefly, compressional waves were insonified perpendicular to the bonded interface by narrow-band transducers. Two types of transducers were employed: a Panametrics 133 A for the ultrasonic inspection during the tensile test, because it is small enough to be mounted in the sample holder of the tensile test stage. The other data discussed in this paper were acquired with a specially made transducer using a PMN-PT (Pb(Mg1/3Nb2/3)O3–PbTiO3) crystal [12]. To provide an almost monochromatic signal, RF pulses of 10 to 30 cycles were generated with a peak

25.2 Experimental and Calibration Procedure

power of up to 3.6 kW at a carrier frequency of 2.25 MHz. The repetition rate was 30 Hz, so that the duty factor was less than 4 ´ 10–4 and therefore the mean RF input power on the transducer was less than 1.5 W. Taking into account the electrical impedance mismatch and the piezoelectric coupling factor of the transmitting transducer leads to an average ultrasonic power of about 0.3 W, thus avoiding heating. The transmitted ultrasonic signals were detected by a broadband transducer (Panametrics V110-22), recorded with a sampling rate of 400 MHz with 8-bit signal depth and Fourier-transformed. In addition to the previous set-up, a phase-sensitive detection unit was added to measure the phase of the ultrasonic signals relative to the RF carrier [13] which excites the transducer. The dependence of the resulting amplitude and phase spectra on the transmitting power was recorded. The transmitting and receiving transducers were coupled to the sample by a thin layer of machine oil and firmly pressed onto the aluminum plates within the tensile test stage. By performing many tests we ensured that this coupling procedure guaranteed linear and reproducible behavior for the measurements discussed here. The evaluation of the data requires knowledge of the absolute stress amplitudes acting on the interface. This makes it necessary to take into account the superposition of all forward and backward traveling waves in the plates bonded together. To obtain an absolute calibration of the wave amplitudes in transmission, a plate of the same material and thickness as that of the composite sample was insonified by compressional waves. For the different frequencies the transmitted displacement amplitude at the free backwall of the plate was measured interferometrically for various excitation voltages of the transducer. Then the same measurements were repeated with the interferometer replaced by the piezoelectric receiving transducer used in the experiments on the bonded samples. This yielded a relationship between the absolute amplitudes of the transmitted waves in the case of a free back wall and the voltage they generated at the receiving transducer. An analogous procedure was carried out with a second plate representing the intromission side. Due to the coupling of the receiving probe the back wall of the plate on the receiving side was no longer free, and neither were the two sides of the plates which were bonded together. This requires further evaluation of the calibrated values to relate the measured results to the forces acting on the interface [10]. The derivation of a relationship between the phases of the different frequencies at the receiving transducer utn and the phases of the force components un at the bonded interface is more complicated. The continuity of stresses and displacements at the aluminum plate/coupling medium and coupling medium/receiver probe interfaces have to be taken into account, which introduces the material parameters of both the coupling medium and the transducer into the calibration equations. The procedure described so far yielded the phases relative to the excitation at the transmitting transducer. To obtain the phases related to the interface vibration, similar considerations exploiting the continuity of stress and displacement at the bonded interface have to be carried out. Details have been presented elsewhere [10] and are omitted here due to lack of space.

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25.3 Calibrated Ultrasonic Transmission Measurements

Samples of two aluminum plates, each 4 mm thick and joined together by a thin epoxy layer about 30 lm thick, were investigated. The aluminum surfaces were polished mechanically to an almost optical finish (root-mean-squared rugosity R = 0.9 lm) and cleaned first with acetone and then with isopropyl alcohol to remove residual impurities. The strain amplitudes eI, e1, e2, and e3, of the incident and the transmitted waves at the excitation frequency and their second and third harmonics were measured and calibrated as described above. At the receiving transducer the phases ut1, ut2, and ut3 of the transmitted waves were determined. So far in our experiments the harmonics higher than the third order have been below the noise level. The standard deviation in the measured strain amplitudes is about 1% for the transmitted fundamental frequency and its second harmonic, and about 5% for the third harmonic. In our experiments transmitted strain amplitudes of up to about 2 ´ 10–4 are obtained. This corresponds to a stress amplitude of about 22 MPa (elastic constant of Al: c11,Al = 107.8 GPa). Because of the stress continuity at the interface, this stress is also present in the bond. This value is close to the tensile strength values of epoxy adhesives [1, 14]. This means that with the ultrasonic strain levels applied here one can cover the part of the stress–strain curve where nonlinear behavior of the bond is expected, while the behavior of the aluminum plates still remains linear. We checked the latter experimentally as well as by estimating the amplitude of the second harmonic using b2 = 7 for aluminum [15]. After a path length of 10 mm the second harmonic generated in aluminum is at least 25 dB below the amplitude of the fundamental frequency. If the distortion of the interface by the higher harmonics generated can be neglected as a higher-order effect, the interface vibration is sinusoidal with the excitation frequency. Its displacement amplitude a0 is determined by the stress and displacement continuity of the waves of fundamental frequency at the interface. If the transmission of the fundamental frequency shows no p hysteresis relative to the interface vibration, we get the result a0 = eBI/k, eBI = 2 e2I e21 . Here k = x/vL is the wavenumber and vL is the compressional sound velocity in the aluminum plates [10]. The strain amplitude in the interface is the ratio of a0 to the thickness of the interface, i.e., the interface strain amplitude is directly proportional to eBI. Figs. 25.1–25.4 show the calibrated measure of the interface strain vibration eBI and the strain amplitudes of the transmitted waves of fundamental frequency and its second and third harmonic e1, e2, and e3 as a function of the incident strain amplitude eI. The solid lines represent a linear fit of eBI and e1, a quadratic fit of e2, and a cubic fit of e3 of the first 12 measuring points up to a strain level eI&1.3 ´ 10–4. The strain amplitudes of the first six measured data points of the third harmonic were below the noise level. The experimentally observed power laws for e1, e2, and e3 render it possible to relate them to an expansion of the stress–strain curve which displays no hysteresis [9]. In this case

25.3 Calibrated Ultrasonic Transmission Measurements Fig. 25.1 Ultrasonic transmission data for a sample of two aluminum plates 4 mm thick bonded together by an adhesive epoxy layer of 30 lm thickness showing the measure of the interface p strain amplitude eBI = 2 e2I e21 versus the input strain amplitude eI and the linear fit of the first 12 measuring points.

Fig. 25.2 Ultrasonic transmission measurement results for a sample of two aluminum plates 4 mm thick bonded together by an adhesive epoxy layer of 30 lm thickness showing the transmitted strain amplitude e1 of the fundamental frequency versus the input strain amplitude eI and the linear fit of the first 12 measuring points.

Fig. 25.3 Ultrasonic transmission measurement results for a sample of two aluminum plates 4 mm thick bonded together by an adhesive epoxy layer of 30 lm thickness showing the strain amplitude e2 of the transmitted second harmonic versus the input strain amplitude eI and the quadratic fit of the first 12 measuring points.

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25 Calibration and Evaluation of Nonlinear Ultrasonic Transmission Measurements Fig. 25.4 Ultrasonic transmission measurement results for a sample of two aluminum plates 4 mm thick bonded together by an adhesive epoxy layer of 30 lm thickness showing the strain amplitude e3 of the transmitted third harmonic versus the input strain amplitude eI and the cubic fit of the first 12 measuring points.

the stress amplitudes caused by the interaction force FIF per unit area in the interface can be calculated by the measured strain amplitudes of the transmitted waves [Eqs. (1 a, 1 b)] [7–10]. FIF c11;Al FIFN

(1 a)

with FIFN e1

2e2 e3

(1 b)

In principle the interaction force curve below the threshold can be determined from the data of only one measuring point [16, 17]. We get FIFN(Da) = (F1N – 3F3N)Da/a12–2F2N(Da/a12)2 + 4F3N(Da/a12)3, where the coefficients contain the measured data of the chosen reference point, here point 12 at eI&1.3 ´ 10–4, i.e., F1N = e1(12), F2N = e2(12), F3N = e3(12). The amplitude of the interface vibration at data point 12 is a12 = a0(12) and Da is the deviation of the interface width from its static equilibrium value. Fig. 25.5 shows the quantity FIFN, which is proportional to the interaction force calculated point by point from the measurements using Eq. (1 b) (solid squares) and determined only by the data of the 12th measuring point (solid line). The phases of the transmitted waves as measured at the receiving transducer are shown in Figs. 25.6–25.8. The phase of the fundamental frequency increases slowly, whereas the phases of the second and third harmonic vary little and the changes are within measuring accuracy, up to data point 12, eI&1.3 ´ 10–4. With further increases of the input strain beyond eI&1.3 ´ 10–4, the interface changes its dynamic behavior because the strain amplitude of the interface vibration is no longer linear with the excitation and the transmitted harmonics change their relative phases. Furthermore their amplitudes no longer follow a power series expansion. We relate this behavior to a hysteresis and increasing viscoelasticity in the interface. The evaluation of the transmitted waves to obtain interface restoring forces has to take the phases explicitly into account. The in-

25.3 Calibrated Ultrasonic Transmission Measurements Fig. 25.5 Ultrasonic transmission measurements for a sample of two aluminum plates 4 mm thick bonded together by an adhesive epoxy layer of 30 lm thickness showing the nonhysteretic approximation of the normalized restoring force [Eq. (1 b)] in the interface calculated point by point from the measurement data (solid squares) and determined only by the data of the 12th measuring point (solid line).

Fig. 25.6 Ultrasonic transmission measurement results for a sample of two aluminum plates 4 mm thick bonded together by an adhesive epoxy layer of 30 lm thickness showing the phase ut1 of the transmitted fundamental wave at the receiver probe versus the input strain amplitude eI.

Fig. 25.7 Ultrasonic transmission measurement results for a sample of two aluminum plates 4 mm thick bonded together by an adhesive epoxy layer of 30 lm thickness showing the phase ut2 of the transmitted second harmonic at the receiver probe versus the input strain amplitude eI.

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25 Calibration and Evaluation of Nonlinear Ultrasonic Transmission Measurements Fig. 25.8 Ultrasonic transmission measurement results for a sample of two aluminum plates 4 mm thick bonded together by an adhesive epoxy layer of 30 lm thickness showing the phase ut3 of the transmitted third harmonic at the receiver probe versus the input strain amplitude eI.

terface vibration amplitude at the threshold is a0&0.2 lm, which corresponds to a strain of about 10–2 for an interface thickness of 30 lm. At strains of this magnitude the transition from elastic to viscoelastic/plastic behavior occurs in quasi-static tensile test experiments [1].

25.4 Ultrasonic Measurement and Destructive Tests

In order to allow both ultrasonic transmission experiments and tensile loading, the ultrasonic set-up was integrated into a small laboratory-scale tensile test stage (Fig. 25.9). The incentive for these experiments originated from the results of combined ultrasonic transmission and dynamic and quasi-static tensile tests of bonded steel tubes [14] and the data shown above. We wanted to test whether there is a correlation of any of the following parameters: the second-order nonlinearity parameter b2, the third-order nonlinearity parameter b3, the distortion factor K, or the interaction force FIF to the destructively determined tensile strengths. The tensile test stage was originally designed to load samples in scanning acoustic microscopes in order to observe in-situ changes in the microstructure of materials [18]. The maximal force which can be applied with this stage is about 5 kN. This limits the cross-section of samples which can be investigated in order to obtain a sufficiently large strain entailing some shear in the bond due to its finite extension. The aluminum plates used to fabricate the test specimens were 5 mm thick in order to increase the stiffness of the plates during tensile loading. Altogether 80 tensile test specimens of aluminum plates 5 mm thick bonded together by an adhesive epoxy layer (Terokal-5070MB-25®1k-EP from Henkel Teroson) 30–50 lm thick with a circular bonding area of 10 mm diameter were fabricated. The two aluminum plates to be joined were first preheated to 50 8C. For 60 of the specimens, one plate was partially waxed in the bonding area with PAT-607/FB in order to produce adhesive defects of different


Fig. 25.9 Schematic drawing of the arrangement of the samples with the ultrasonic transmitting and receiving transducers and their integration into a tensile test stage with the location of the interface.

diameters. The epoxy adhesive was applied to the other plate. The plates were pressed against each other gently. The ensemble was cured for 45 min at 180 8C and allowed to rest for 24 h. In order to avoid an undefined support in the test stage, the epoxy that had escaped from the joint during the heat treatment was carefully removed from the perimeter of the bonding area with an abrasive wheel. Fig. 25.10 shows one of the samples after the bond failed in the tensile test.

Fig. 25.10 A bonded interface of a tensile test specimen after failure. The bond, diameter 10 mm, contained an adhesive defect (diameter 3.5 mm) obtained by waxing. The adhesive was Terokal-5070MB-25®1k-EP from Henkel Teroson.

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Before the tensile test the samples were investigated by ultrasonic transmission measurements as described in Section 25.2. The peak power of the RF-carrier pulse (again 10–30 cycles, center frequency 2.25 MHz) was swept from 0 up to 3.6 kW and back to zero. The transmitted ultrasonic signal was detected by a broadband receiver probe, recorded, and Fourier-transformed. The dependence of the resulting amplitude and phase spectra on the transmitting pulse power was recorded. Figs. 25.11 and 25.12 show the results obtained for two of the specimens, one with a weak and one with a strong bond of 5.5 and 32.5 MPa tensile strength, respectively. The tensile strength of the bonds was determined later by loading the aluminum samples until fracture and dividing the measured force by the nominal area of the bond, here 78.5 mm2. The loading procedure was displacement-rate controlled (0.5 lm s–1), resulting in a strain rate of e_ &10–2 s–1. The amplitudes of the transmitted waves of fundamental frequency (Fig. 25.11) and of the second and the third harmonic (Fig. 25.12) are plotted in arbitrary units as recorded by the receiving transducer as a function of RF-input peak power of the transmitting probe. In the case of large adhesive defects within the bond, the transmitted amplitudes of all signals show hysteresis as a function of input RF power. For small defects or defect-free bonds, no hysteresis could be observed. After the ultrasonic measurements the specimens were loaded until fracture to obtain the tensile strength. During the loading procedure nonlinear ultrasonic transmission measurements with an excitation peak power of 1.86 kW were carried out. Figs. 25.13 and 25.14 show the results. The amplitudes of the transmitted waves of fundamental frequency (Fig. 25.13) and of the second and the third harmonic (Fig. 25.14) are plotted in arbitrary units as recorded by the receiver probe. The horizontal axis represents the number of measuring points

Fig. 25.11 Transmitted amplitudes for a sample of two aluminum plates 5 mm thick bonded together by an adhesive epoxy layer of 30–50 lm thickness. The amplitudes of the fundamental frequency transmitted through (a) a weak bond (tensile strength

5.5 MPa) and (b) a strong bond (tensile strength 32.5 MPa) are plotted versus the input RF pulse power. The amplitudes show hysteresis for the weak but not for the strong bond.


Fig. 25.12 Transmitted amplitudes for a sample of two aluminum plates 5 mm thick bonded together by an adhesive epoxy layer of 30–50 lm thickness. The amplitudes of the second and the third harmonic generated in transmission through (a) a weak

bond (tensile strength 5.5 MPa) and (b) a strong bond (tensile strength 32.5 MPa) are plotted versus the input RF pulse power. The amplitudes show hysteresis for the weak but not for the strong bond.

Fig. 25.13 Ultrasonic transmission measurement results (excitation RF peak power 1.86 kW) carried out on samples of two aluminum plates 5 mm thick bonded together by an adhesive epoxy layer of 30–50 lm thickness during quasi-static tension loading (displacement rate 0.5 lm s–1) until fracture. The amplitudes of the fundamental frequency transmitted through (a) a weak

bond (tensile strength 5.5 MPa) and (b) a strong bond (tensile strength 32.5 MPa) are plotted in arbitrary units as recorded by the receiving transducer. The horizontal axis represents the number of measuring points. Three data points were measured per second. Fracture occurred at a loading force of 432 and 2550 N, respectively.

and hence the loading force. Due to the backlash in the tensile test stage at the beginning of a test, it is difficult to designate units of force on the abscissa in Figs. 25.13 and 25.14. As stated above, the loading procedure was displacementrate controlled (0.5 lm s–1). This entailed strain rates that varied for samples

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Fig. 25.14 Ultrasonic transmission measurement results (excitation RF peak power 1.86 kW) carried out on samples of two aluminum plates 5 mm thick bonded together by an adhesive epoxy layer of 30–50 lm thickness during quasi-static tension loading (displacement rate 0.5 lm s–1) until fracture. The amplitudes of the second and the third

harmonic wave generated in transmission through (a) a weak bond (tensile strength 5.5 MPa) and (b) a strong bond (tensile strength 32.5 MPa) are plotted in arbitrary units as recorded by the receiving transducer. The horizontal axis represents the number of measuring points. Three data points were measured per second.

which differed in interface thickness, while the stress rates varied not only with the interface thickness, but also with differences in elastic and viscoelastic properties or in any combination of these parameters. Three measuring points were taken per second. Fracture occurred at F = 432 N for the weak bond, which corresponds to a nominal stress of r = F/A = 5.5 MPa if one divides F by the total area A of the bond (A = 78.5 mm2). For the strong bond fracture occurred at F = 2550 N corresponding to r = 32.5 MPa. The defects were delaminations with contacting surfaces (‘kissing bonds’) which weaken the overall bond strength by reduction of the adhesive strength in the waxed area. This explains the high ultrasonic transmission through the weak bond. The almost unhindered ultrasonic transmission through “kissing bonds” which are under compressive stress is very well known in nondestructive testing. It has been shown recently [14, 19, 20] that it is indeed possible to open and close kissing bonds reversibly with tensile or compressive stress, so that they become reflective and show a contrast in ultrasonic C-scan images if opened. For both samples the transmitted amplitude of fundamental frequency first increases slightly with the static tensile stress, which might be caused by a stiffening of the adhesive layer under ultrasonic strain; then the amplitude decreases again, which might indicate a softening, i.e., degradation of the adhesive bond, until finally fracture occurs. The effect is much stronger for samples with large adhesive defects than for defect-free samples or samples with small adhesive defects. The effect increases as the size of the defective area increases. The amplitudes of the higher harmonics (Fig. 25.14) generated in transmission


through the strong bond show a similar behavior to the amplitude of the fundamental frequency. The slight increase of the higher harmonics at the beginning might be caused by the slight increase in the overall transmission coefficient. The amplitude variations of the higher harmonics generated by the weak bond are much more diversified, which probably accounts for more drastic changes in the dynamic behavior (linear and nonlinear elastic, viscoelastic, plastic, generation and growth of delaminations, etc.) of the interface under increasing quasi-static load. Comparing the size and shape of the defects observed in the bonding areas after fracture with the variation of the amplitudes of the higher harmonics as a function of load, let us conclude that there is a strong interdependence. The variation of the generation of the higher harmonics with the quasi-static load is strong (Fig. 25.14) if the defective area is perfectly round and sharply separated from the rest of the bonding area (Fig. 25.10). If the perimeter of the defective area is serrated, the change of the higher harmonics with load is barely noticeable. In Section 25.3 we based our analysis of the data shown in Figs. 25.1–25.8 on the assumption that no hysteresis occurs in the interface vibration up to an input strain level of eI&1.3 ´ 10–4. The abrupt change in the phases of the transmitted waves with increasing excitation amplitude indicates the onset of interface vibration hysteresis. The hysteresis shown in Figs. 25.11 and 25.12 might, however, be caused by a spatial rearrangement of the units of the microstructure of the epoxy polymer as a function of applied ultrasonic strain. Such a mechanism has been invoked in explaining the linear frequency dependence of the ultrasonic absorption mechanism observed in many polymers [21]. Other groups contested this idea and explained this absorptive behavior by a thermally activated structural relaxation mechanism with a broad distribution of relaxation times [22]. In this context we should like to mention the structural units in amorphous materials, including polymers and adhesive resins which exhibit relaxation over very large time scales, from nanoseconds to hours [23, 24]. These relaxation centers manifest themselves on the one hand in the tunneling phenomena observed at very low temperature [25], and on the other hand in the thermally activated motion of structural units responsible for ultrasonic absorption and internal friction. In another context the spatial rearrangement of the microstructure of a material with applied dynamic strain has been termed “conditioning” [26]. For example, the stress slowly relaxes when a strain is applied to a rock containing micro-cracks, because the micro-cracks accommodate by changing their opening displacement [26–28]. Likewise such effects, caused by dislocation movements over limited displacements induced by ultrasonic strains, have been observed in metals [29]. The tensile strengths of the many samples examined are related to different ultrasonic measurement quantities. Fig. 25.15 a shows the distortion factor K as a function of tensile strength with the transmitting transducer excited at an RF peak power of 3.6 kW. The nonlinearity parameter b2 ! A2/A21 (Fig. 25.15 b) informs about the second-harmonic, and the nonlinearity parameter b3 ! A3/A31 (Fig. 25.15 c) about the third-harmonic generation. These three ultrasonic quan-

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Fig. 25.15 (a) Distortion factor K; (b) parameter A2/A21; (c) parameter A3/A31 derived from the ultrasonic transmission data at an excitation RF peak power of 3.6 kW versus tensile strength. The samples consisted of two aluminum plates 5 mm thick bonded together by an adhesive epoxy layer of 30–50 lm thickness.

tities show little if any correlation to the tensile strengths, even if we account for large error bars in the tensile strength measurement results because of bending effects in the aluminum plates by the load in the tensile test apparatus. Furthermore, there is still a certain influence on the data by the higher harmonics generated by the Panametrics transmitting transducer which is difficult to quantify in situ. Tensile loading of samples of two bonded 5 mm-thick aluminum plates still caused bending, although much less than in case of the bonded plates that were 4 mm thick. Fig. 25.16 shows (a) the phase ut1 of the transmitted wave of fundamental frequency measured at the receiver probe and (b) the interaction force amplitude per area [7–10], using Eq. (1) versus tensile strength. As before, the excitation peak was at 3.6 kW. In contrast to the results shown in Fig. 25.15, there is a tendency for the phase of the transmitted fundamental wave (Fig. 25.16 a) and the interaction force amplitude (Fig. 25.16 b) to increase with increasing tensile strength. The increase in the phase of the transmitted fundamental wave is expected theoretically by evaluating the boundary conditions at the transmitting and the receiving transducers as well as at the bonded interface, taking into account the forward and backward propagating waves in the plates. If we assume


Fig. 25.16 (a) Phase of the transmitted wave of fundamental frequency recorded at the receiving transducer, and (b) restoring force FIF derived from the transmission measurement data at an excitation peak power of

3.6 kW versus tensile strength. The samples consisted of two aluminum plates 5 mm thick bonded together by an adhesive epoxy layer of 30–50 lm thickness.

no hysteresis in the interface vibration (i.e., the interface thickness and the corresponding interaction force vibration are in phase) for the full range from a complete delamination to a perfect bond, the theoretically expected phase shift is smaller (see Fig. 25.17 a) than the measured effect. If we assume that there is an increase in the hysteretic behavior of the interface vibration with decreasing bond quality, the magnitude of the measured effect can indeed be obtained also in the calculations (see Fig. 25.17 b). The increase in the hysteretic behavior

Fig. 25.17 Calculated phase of the transmitted wave of fundamental frequency as a function of the ultrasonic transmission coefficient of the interface in the range from 0 (complete delamination) to 1 (perfect bond, i.e., complete transfer of ultrasound): (a) without hysteresis; (b) with interface

vibration hysteresis which increases from 0 to –p/2 with the decrease in the ultrasonic transmission coefficient from 1 to 0. The hysteretic phase shift in (b) almost encompasses the phase shift observed experimentally (Fig. 25.16 a).

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means a decrease in the phase of the transmitted wave, i.e., the waves are transmitted more slowly through the interface with decreasing tensile strength. The theoretically expected interaction force calculated from the measured ultrasonic data at an excitation peak power of 3.6 kW corresponds to the measured tensile strength, as can be seen in Fig. 25.16 b. There are several reasons for the scatter of the data: (a) nonlinear contributions of the transducers and measurement equipment are superimposed on the measurement effect; (b) the samples bend during a test, producing both tensile and shear stress in the bonded interface; and (c) there are expected to be statistical variations of the properties of the epoxy bonds themselves, due to different microstructures and curing. Furthermore, depending on the dynamic behavior of the interfaces, the deviation of the phases of the transmitted waves from their values in the nonhysteretic case must be taken into account in the evaluation of the interaction forces, especially if the ultrasonic excitation is high enough to obtain ultrasonic stresses comparable with the bond strength.

25.5 Conclusion

Samples consisting of aluminum plates joined together by thin epoxy adhesive layers were investigated by nonlinear ultrasonic transmission measurements. A threshold behavior of the transmitted harmonics was observed. Their amplitudes depend on the excitation following the power series expansion of a quasistatic interaction force and the phases vary little at low-amplitude excitation. Exceeding the threshold causes a considerable change in the dynamic behavior of the interface. The description of thin-bonded interfaces only by binding forces, without taking the material properties of the adhesive layer explicitly into account, has been used to determine interaction forces in the interface by calibrating the strain amplitudes of the ultrasonic wave. The nonlinear ultrasonic transmission measurement data were compared with tensile test results. The phase of the transmitted wave at the excitation frequency and the interaction forces in the interface correlate with the tensile strength. These might be used as nondestructive testing parameters.

Acknowledgments

We the GE the

benefited from the European Science Foundation program NATEMIS and European Union project FP6-502927 AERONEWS. We thank Dr. G. Splitt, Inspection Technology Systems, Cologne–Hürth, Germany, for providing PMN-PT transducers.

References

References 1 See, for example, G. Habenicht, Kleben, 2 3

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3rd edn., Springer, Berlin, 1997. D. C. Hurley, C. M. Fortunko, Meas. Sci. Technol. 1997, 8, 634–642. S. U. Faßbender, W. Arnold, Rev. Progr. QNDE 1995 (Eds.: D. O. Thompson, D. E. Chimenti), Plenum Press, New York, 1996, Vol. 15, pp. 1321–1328. O. Buck, W. L. Morris, J. M. Richardson, Appl. Phys. Lett. 1978, 33, 371–373. M. Rothenfußer, M. Mayr, J. Baumann, Ultrasonics 2000, 38, 322–326. C. Bockenheimer, D. Fata, W. Possart, M. Rothenfußer, U. Netzelmann, H. Schäfer, Int. J. Adhes. Adhesives 2002, 22, 227–233. S. Hirsekorn, A. Koka, A. Wegner, W. Arnold, Rev. Progr. QNDE 1999 (Eds.: D. O. Thompson and D. E. Chimenti), Plenum Press, New York, 2000, Vol. 19B, pp. 1367–1374. A. Wegner, A. Koka, K. Jansen, U. Netzelmann, S. Hirsekorn, W. Arnold, Ultrasonics 2000, 38, 316–321. S. Hirsekorn, Ultrasonics 2001, 39, 57– 68. S. Hirsekorn, A. Koka, W. Arnold, 2nd Workshop “NDT in Progress”, Prague, Czech Republic, Oct. 6–8, 2003, pp. 99– 106. M. Schlimmer, University of Kassel, private communication, 2004. TRS Ceramics Inc., State College, PA 16801, USA, www.trsceramics.com Ritec, Advanced Measurement System Model RAM-5000-SNAP, www.ritecinc.com A. Koka, S. Hirsekorn, W. Arnold, R. Hunke, J. Häberle, M. Schlimmer, “Zerstörungsfreie Detektion von Klebverbindungsfehlern mit Ultraschall und Untersuchung der Auswirkung dieser Fehler auf die mechanische Beanspruchbarkeit der Verbindung”, IZFP Report No. 040403-E, 2004, and “Schriftenreihe des Instituts für Werkstofftechnik”, Universität Kassel, Report No. 3-2004, 2004. M. A. Breazeale, J. Philip, in Phys. Acoustics, Vol. 17 (Eds.: W. P. Mason, R. N.

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26 Debonding of Pressure-Sensitive Adhesives: A Combined Tack and Ultra-Small Angle X-Ray Scattering Study E. Maurer, S. Loi, and P. Müller-Buschbaum

Abstract

The debonding of pressure-sensitive adhesives (PSAs) is investigated by a combination of the mechanical tack test and ultra-small angle X-ray scattering (USAX). Two chemically different PSAs are investigated, namely the homopolymer poly(n-butyl acrylate) (PnBA) and a statistical copolymer with a monomer ratio of 80% 2-ethylhexyl acrylate (EHA) to 20% methyl methacrylate (MMA). Scattering yields additional structural information on a nanoscopic level, which is identified as nano-bubbles located inside the PSA film. Due to their size, these nano-bubbles are invisible with optical techniques. The probed nano-bubbles, as a substructure of vertically expanded adhesive material, show that the structure creation during debonding, which is well known at macroscopic length scales, also proceeds in the nanometer range.

26.1 Introduction

Pressure-sensitive adhesives (PSAs) [1, 2] can be stuck to a huge variety of materials by applying a slight pressure. Based on the behavior of their polymer melts, the rheological properties of PSA play a key role in the phenomenological appearance of tackiness [3–6]. From the mechanical point of view, a PSA polymer melt possesses a viscoelastic nature. At long time scales the viscous aspect dominates and enables the adhesive to achieve intimate contact with the surface of a solid specimen, adapting to the latter’s surface profile [7, 8]. On the other hand the elastic properties, which gain mainly in importance at short time scales, allow the adhesive to sustain short-term shear forces. The fact that PSAs combine both aspects at the same time distinguishes them from other types of adhesives that are converted from a fluid state to a solid state, for example, by a change in temperature or through a chemical reaction. The lack of hardening in PSA application enables a desired, controlled release of the adhesive bond. One of the most prominent exAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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26 Debonding of Pressure-Sensitive Adhesives

amples in daily life is the stick-on note, in addition making use of the ability to undergo several cycles of bonding and detaching. Scientifically, the quality of adhesion can be quantified in the so-called tack test [9–12]. A probe punch, like a flatended, rigid cylinder, is pressed with a defined force onto a PSA film [13]. After the force has been maintained for a well-defined time, the punch is withdrawn from the surface with a fixed velocity. During the whole process the force needed to sustain the constant retraction velocity is probed as a function of the distance between the film and the punch. Usually, the resulting curves show characteristic features such as a sharp maximum followed by an extended force plateau. Even at film– punch distances of up to multiple film thicknesses, a nonzero force value is detected. The geometry of the tack test with a flat-ended cylinder [14] ensures, contrary to a comparative test with a spherical indenter [15], a uniform spacing of substrate and punch [16]. Despite the homogeneous elongation throughout the contact area, a highly heterogeneous structure of cavities and fibrils develops in the polymer [17]. In other words, the material has to face the challenge to occupy a rapidly increasing volume and to respect its low compressibility as well. The debonding process during the tack test was analyzed in detail by optical microscopy both from underneath, through a transparent substrate, and from the side. So far, four different characteristic stages of adhesive failure have been discovered by means of optical techniques (see Fig. 26.1) [11, 17–19]. Firstly homogeneous elongation of the polymer film in the direction of tension is accompanied by a sharp increase in the force. As traction proceeds, cavities are introduced into the film locally [3, 20–24]. With an accompanying stress release, the appearance of cavities corresponds to the force maximum. As the cavities expand laterally, the force promoting debonding suddenly decays to a comparatively low, but nonzero, value. A plateau of constant height in the force versus distance curve follows and the cavities already occupying most of the nominal contact area now expand, mainly vertically.

Fig. 26.1 According to Creton and Lakrout [18] four stages of debonding, each represented by a sketch in the top row, are assigned to a corresponding part (broken line) of the force–distance (F–d) curve. During all stages the punch is retracted at velocity v from the substrate. The evolution of cavities

embedded in the polymeric film is crucial for the classification. The stages are denoted following the temporal development of the material–cavity ensemble as homogeneous expansion, appearance of cavities, lateral expansion, and vertical expansion.

26.2 In-Situ Small Angle Scattering Using Synchrotron Radiation

As the cavities remain well separated the polymeric material between them still provides a local connection between the punch and the substrate. As debonding proceeds, the film is transformed out of a foam-like state to a fibrillar structure. Finally, when air rushes in from the outer border of the vertically expanded film, the measured force vanishes to zero. Thanks to investigations of PSA failure by optical microscopy, a detailed understanding of the highly nonlinear force–distance curve in the test has been obtained. Nevertheless all experiments are restricted toward small length scales by the optical resolution limit. In order to overcome this limit we applied X-ray scattering techniques [25]. As a further advantage, scattering is bulk-sensitive, meaning that an insight into internal structures is possible.


Small-angle X-ray scattering (SAXS) is a standard testing method in materials science [26, 27]. It provides statistically significant information about scattering density heterogeneities in the nanometer range. Among the numerous applications for SAXS, the determination of particle sizes in colloidal systems is a prominent example [28]. SAXS experiments have also been combined with simultaneous mechanical testing in order to probe structural stress response [29– 32]. The idea of studying density heterogeneities can be also applied to adhesives. In the tack test, during debonding the formerly homogeneous adhesive film changes to a foam-like state. Density heterogeneities are introduced into the system on several length scales. With SAXS, the hitherto explored regime of length scales is enlarged toward the molecular range [25]. In matter exposed to X-rays, electrons are excited by the incoming radiation. In the case of X-rays the wavelength is of the order of 1 Å. Thus the volume elements filled with polymer material scatter more intensively than regions of vacancies. The scattered intensity is measured as a function of the scattering angle h, which is between the direction of incidence i and the analyzed direction f. During the scattering process X-ray photons undergo a momentum transfer ! ! ! q kf ki , depending on the scattering angle. As energy is conserved, the absolute value of the momentum transfer can be expressed as Eq. (1). qj j!

4p h sin k 2

1

Summing all the scattering amplitudes of the emitted waves and respecting their correct phases leads to the total scattering amplitude given by Eq. (2), where q! r is the electronic density representing the spatial distribution of electrons in the sample [26, 27].

q A!

Z

d3 rq! r e

!!

iqr

2

423

424


Formally A ! q is the Fourier transform of the electron density. In the region of small ! q, corresponding to small scattering angles h, it is dominated by low spatial frequencies in the electronic density. Depending on the probed regime of length scales D, the descriptions “small-angle scattering” (SAXS) and “ultrasmall-angle scattering” (USAX) have established [33]. The scattered intensity, which is proportional to the signal registered at a detector, is given by the absolute square of the amplitude. A peak in the intensity for a certain momentum transfer ! qmax represents a periodicity in the correlation function of the electron density and a frequent presence of the corresponding spacing D, with D 2p=j! qmax j. For in-situ SAXS experiments during the tack test, the use of a synchrotron source providing highly intense X-ray radiation is necessary. The experiments presented in this paper were performed at HASYLAB/DESY at the beamline BW4. This beamline enables SAXS and USAX experiments in the common transmission geometry with a quasi-monochromatic beam, operating a two-dimensional position-sensitive detector. The direct beam has to be shadowed with a beam stop, because its extremely high intensity cannot be handled by the detector. The scattered intensity containing the desired information is orders of magnitude smaller. To enhance the sensitivity with respect to small intensities, the flight path is under high vacuum, suppressing air-scattering perturbation. The sample itself and the surrounding tack apparatus are in air. To fulfill all the necessary boundary conditions related to a scattering experiment at a synchrotron beamline, a special tack apparatus was constructed. Firstly, its geometry did not disturb the in- and outgoing beams. Furthermore, its size and weight allowed convenient goniometer handling at the beamline. The sample, a PSA film (with a thickness between 10 and 100 lm, on top of a float glass slide), was placed parallel to the incident beam, shadowing only a small part of the beam (approximately 400 lm in diameter). While film and beam both remained fixed in real space, the tack test was performed in such a way that the punch was retracted from the PSA in the direction perpendicular to both film surface and incoming beam. Thus during debonding the polymer material between the film surface and the punch was illuminated and probed. The two-dimensional position-sensitive detector was mounted perpendicular to the incoming beam. The distance from the detector to the punch was 13 m, corresponding to a USAX experiment. At an X-ray wavelength of 0.138 nm heterogeneities in the range 25–400 nm were resolvable. The upper limit was due to geometric constraints introduced by the size of the beam stop. Toward small structures the resolution was limited by the finite size of the detector. Scattering patterns were recorded at several film–punch distances. To clarify the geometry, we define the x-axis as the direction of the incident beam. Thus the film was located in the (x, y)-plane whereas traction in the tack test was directed along the z-axis. In Fig. 26.2 the geometry of the scattering experiment is shown schematically, together with a typical scattering signal recorded by the detector. The contribution of each component in the set-up to the two-dimensional scattering image is identified by a stepwise approach to the final experimental condi-


Fig. 26.2 (a) In the scattering geometry the volume underneath the probe punch is illuminated from the side by the X-ray beam. The incoming beam is placed parallel to the film surface. The film and substrate form shadows in only a small part of the beam. Note that the relative dimensions of the punch, beam, and film thickness are realistic as depicted, whereas the lateral extension of

the film and the substrate dimensions are less than the realistic values. (b) A typical signal on the position-sensitive, two-dimensional detector. The dimensions of the detector are small compared with its distance from the sample, meaning that only radiation scattered in a range of small angles (below 0.38) is registered.

tions (see Fig. 26.3). Most of the primary beam is blocked by a beam stop. Nevertheless the position of the primary beam is estimated by the radial symmetric decay of intensity surrounding the beam stop shadow. As the film surface is placed in the beam the incoming radiation grazes the surface. Evanescent waves thus enter the film; they are the source of the diffuse scattering registered above the direct beam position. As the secondary radiation is emitted in a relatively wide angular distribution in z-direction as compared to the horizontal divergence of the beam, the diffuse scattering is shaped as a vertical streak on the detector. Strong horizontal streaks (in the y-direction) on both sides of the direct beam position appear only after the tack test has been performed. Resulting from the mechanical test, these streaks are thus the signal of primary interest. As the USAX signal is not isotropic but is concentrated in horizontal streaks, qualitative conclusions about the form and orientation of objects composing the polymeric material beneath the punch can already be made. Firstly, anisometry of the USAXS signal reveals well-orientated objects. Furthermore, the comparatively small extension of these streaks in the vertical direction argues for highly elongated objects showing a large aspect ratio [34]. Remembering the geometry of the tack experiment and keeping in mind the macroscopically observable fibrils, both conclusions appear very reasonable. For a more qualitative evaluation of the data, we perform a horizontal cut of the two-dimensional intensity, restricted to the horizontal momentum transfer qz 0. In order to improve the counting statistics, a slab of intensity is integrated in the perpendicular direction.

425

426


Fig. 26.3 (a) Monitoring of the direct beam. A circular incision in the intensity pattern, due to the beam stop mounting, is observed. (b) The presence of the polymer film in the beam causes diffuse scattering on the film surface, resulting in an enhanced scattering intensity in the vertical direction. (c) Next the punch is brought close to the

film but still remains out of contact. Diffuse scattering from the punch surface results in an additional vertical streak in the lower part of the scattering image. (d) Not until the tack experiment is performed does a clearly pronounced horizontal streak appear, containing information on the microscopic nature of the tack test.

26.3 Microscopically Inaccessible Substructures

We investigated two PSA model systems, namely the homopolymer poly(n-butyl acrylate) (PnBA) [25, 35–38] and a statistical copolymer with a monomer ratio of 80% 2-ethylhexyl acrylate (EHA) to 20% methyl methacrylate (MMA) [39]. The molecular weight Mw of PnBA is centered around 260 kDa with a polydispersity of Mw/MN = 3.78. The statistical copolymer (EHA-co-MMA) is a basic component of industrial PSA with an accordingly broad molecular weight distribution. PSA films were prepared by solution casting on top of float glass slides, which had been pre-cleaned in an acid bath before preparation [40]. The cleaning bath consisted of 100 mL 80% H2SO4, 35 mL H2O2 and 15 mL deionized water. After 15 min at 80 8C in the acid bath, the substrates were taken out, rinsed in deionized water, and dried with compressed nitrogen. The polymers were dissolved in toluene (PnBA) or isobutanol (EHA-co-MMA), respectively. Each solution was distributed on top of horizontally placed slides. By controlled solvent evaporation, films of homogeneous thickness between 10 and 100 lm, depending on the concentration of the solution, remained. The tack test was performed in fixed experimental conditions: the probe (a glass cylinder of 2 mm diameter) was pressed with a force of 6 N onto the PSA film surface. After 10 s of contact time the probe was withdrawn with a constant velocity of 1050 lm s–1. In Fig. 26.4, horizontal cuts for both sample systems are displayed. The curves correspond to different distances between the


punch and the film surface. For clarity the scattering curves have been shifted upward along the logarithmic intensity axis with increasing film–punch distance. Comparing the homopolymer and the statistical copolymer, the corresponding scattering data show qualitatively different features. Whereas the curves measured on PnBA show one strong peak (I) and a subsequent broader one (II), the data for the copolymer exhibit multiple, clearly pronounced fringes. In addition, the cuts on the PnBA data are symmetrical with respect to the position of the direct beam, contrary to those on the P(EHA-co-MMA). The symmetry of the first data set makes it possible to improve the statistics by adding the signal for both signs of momentum transfer. Consequently, the curves are displayed versus the absolute momentum transfer. This procedure is clearly not possible in the case of P(EHA-co-MMA). Both the intensity and the angular positions of the fringes depend on the sign of the momentum transfer. We will refer to the latter aspect in the following discussion as positional asymmetry. Fig. 26.4 furthermore shows that the respective features for both PSA systems become less pronounced with increasing film–punch distance. The broad peak II in the PnBA data vanishes. For the statistical copolymer, the fringes decrease in amplitude and the degree of asymmetry diminishes. In view of the obviously deviating features of the resulting curves, the data fitting was performed in different ways. As nonhomogeneous decay of the scattering intensity was observed for both polymers, a substructure beyond the optical resolution has to exist in both systems. Joining Gay and Leibler’s [41] suggested picture of micro-bubbles exist-

Fig. 26.4 Horizontal cuts corresponding to different film–punch distances for (a) PnBA and (b) P(EHA-co-MMA). The symbols represent measured data whereas the solid lines depict fits to the data. The broken lines in both plots mark the resolution limit for small momentum transfers due to the finite

size of the beam stop. For clarity all curves are shifted along the intensity axis. From the bottom to the top the film–punch distance increases stepwise in both plots for (a) by 125 lm starting at 125 lm, and for (b) by 50 lm starting at 100 lm.

427

428


ing within the macroscopic fibrils, we can imagine this substructure as an assembly of empty cavities embedded in the polymeric material, but contrary to those authors we denote these cavities as nano-bubbles or nano-cavities in order to match the length scales addressed by the scattering experiment. Because the scattering data are dominated by the shape and distribution of the optically inaccessible nano-objects, the differences in the intensity observed has to result from the difference in the type of PSA under investigation, due either to chemical differences or perhaps more probably due to differences in the rheological properties of PnBA and P(EHA-co-MMA). In terms of the stability of the nanocavities filled with vacuum, the nonequilibrium character of the debonding process should be kept in mind. First we focus on the interpretation of the data on the homopolymer as plotted in Fig. 26.4a [25]. Following the geometry of the tack test, we assumed the nano-bubbles to be cylindrical in shape with the footprint in the (x, y) plane. In detail, the mean radius of these cavities should be denoted by r. The height of the cylinders remains unspecified, as it does not enter the horizontal cut performed. We are restricted by the qualitative statement, already concluded above from the small vertical extension of the streaks, that the height should greatly exceed the cylinder radius. Furthermore the spatial distribution of the cylindrical cavities is supposed to show a statistically significant, highly prominent, next-neighbor distance between the nano-cavities. Regarding the physics of the scattering experiment, the electron density in real space of the probed volume can be expressed as a convolution of the generalized cavity-form cylinder, and a structural function reflecting the location of each cylindrical object. In reciprocal space the information on form and structure factorizes to a product of the form factor contribution F~ q, the structure factor contribution S~ q and the diffuseness of the interface R~ q [Eq. (3)]. I~ q N S! q F! q R! q B

3

The constant background is denoted B. The number of scattering objects N is adapted to the overall scattering intensity. For fitting, the scattering data have to be compared with the appropriate one-dimensional cut I! q qx 0; qy ; qz 0 out of the three-dimensional function I! q . In general, wave vectors that are being only slightly deflected from the incoming direction differ in momentum from the incoming radiation mainly in directions perpendicular to the incoming beam. Thus in small angle scattering the momentum transfer in the qx direction is often negligible, whereas qz is set at zero due to the particular way the cut from the measured scattering pattern was performed. Introducing the assumed model of an assembly of cylindrically shaped cavities, the form factor F ! q is expressed by the absolute square of the 3D Fourier transform of a cylindrical object (the cavity), and its respective one-dimensional cut is given by Eq. (4), where J1 denotes the Bessel function of the first kind and order 1.


F0 qy ; r r 4

J1 qy r 2 qy r

4

As the radii of the cylinders are most probably polydisperse, the form factor F0 qy was folded with a log normal distribution f r of the cylinder radii r [Eq. (5)]. Z1 Fqy

F0 qy ; r f r dr

5

0

Furthermore, the structure factor can be expressed within a paracrystalline model by Eq. (6) [42]. Sqy

1 1

fq qy

2fq qy cos qy D fq qy 2

with

fq qy exp 2q2y r2

6

A local deviation of the distance between next-neighboring cylinders from its mean value is quantified by the roughness parameter r. The diffuseness of the interfaces is modeled with a simple roughness parameter rR [Eq. (7)] [25]. Rqy exp q2y r2R :

7

All fits to the scattering data utilizing the model described are shown by the solid lines in Fig. 26.4 a. The fit shows a perfect agreement with the data for large momentum transfers, including the broad peak II [25]. A good qualitative accordance in the peak shape and its position could be achieved for low momentum transfers around the peak denoted I. For all cuts, peak I is present and at a constant position in qy, revealing a most prominent average spacing of the cylinders, with a fixed value of D = 300 ± 20 nm. The appearance of peak I in all the curves proves that the substructure is present in real space during the entire debonding process. As long as peak II is detectable in the data, it also remains at a fixed qy. Therefore a constant value of the average cylinder radius can be stated, namely r = 98 ± 5 nm. The decay of both peak intensities with increasing film–punch distance is interpreted in terms of a subsequent loss of nano-bubbles. This idea follows Gay and Leibler’s suggestion [41] that with increasing tension the micro-bubbles merge and thus decrease in number [25]. Obviously, the deviating features of the scattering data for the statistical copolymer P(EHA-co-MMA) have to be interpreted in a different way [39]. Each fit has to respect the most eye-catching feature of the curves, namely the positional asymmetry. Thus the first step is to quantify this asymmetry. For this purpose the minima in intensity were numbered by the index j and plotted versus the corresponding scattering angle h. Negative j values correspond to negative scat-

429

430


tering angles. A parabolic dependence is observed which is exemplified in Fig. 26.5 for the scattering curve measured at a film–punch distance of 100 lm. Focusing first on aspects of the scattering experiment, one simple model to explain positional asymmetry of the kind observed is scattering at a one-dimensional grating of small objects, e.g., voids, that is tilted in the (x,y)-plane [39]. Geometrically this means that each void, displaced with respect to the neighboring one by Dy0 in the y-direction, is shifted additionally by Dx in the beam direction (see Fig. 26.5). Basic geometrical considerations for the path difference DS between neighboring voids lead to a small angle approximation to a parabolic correlation of the index j with the corresponding scattering angle h (Eq. (8), where n refers to the total number of object and k is the X-ray wavelength). jh

n h2 Dy0 h Dx 2 k

8

Inversion symmetry in j (h) is broken by Dx h2/2. Thus the ratio g = Dx/(2 Dy0) is considered as a parameter governing the degree of positional asymmetry. In detail 42 < g < 50 holds for all the scattering curves in Fig. 26.4 b. The total number of objects n cannot be determined by the parabolic fit. Nevertheless, a lower limit can be extracted from the number of side minima i (fringes). As 11 minima have been observed on one side in the scattering curves, n has to exceed 2i = 22. For all film– punch distances the relation 210 nm < nDy0 < 257 nm is valid. In the actual case of PSA debonding, the physical origin of the scattering signal might again be nano-bubbles embedded in the polymer material. Princi-

Fig. 26.5 (a) An index j was assigned to each minimum in scattering intensity and plotted versus its angular position. A parabolic fit (solid line) was performed. (b) A grating of voids tilted in the (x, y)plane. Each void is shifted with respect to the neighboring one by Dy0 in the y-direction

and in addition by Dx in the x-direction. Representative X-ray traces (broken lines), entering along the i-direction and scattered at an angle h, for example, are shown for clarity. In the leftmost void, the path difference Ds with respect to the neighboring ray is indicated by a dotted line.


pally, the above proposed model explains positional asymmetry in terms of a structural arrangement of objects such as voids or elongated nano-bubbles rather than in terms of the object shape. Taking up this idea and transferring it to the context of adhesives, the scattering signal may be due to a regular arrangement of nano-bubbles. Contrary to the evaluation of the data relating to PnBA, the shape of the nano-bubbles remains unspecified in detail. Furthermore, with respect to the one-dimensional character of the assumed model, the most likely counterpart in real space seems to be the partition walls separating two neighboring cavities (see Fig. 26.6). Thus we think that the polymer partition walls contain a linear, regular arrangement of nano-bubbles whose projection in the (x, y)-plane forms a tilted grating in the geometry of Fig. 26.5 b. Note that the proposed interpretation of the scattering data might not be unique. Further tests to clarify the geometry in real space are indispensable. Once a model for positional asymmetry has been established, in other words the phase of the scattered waves has been described quantitatively, the scattering data have to be fitted with respect to intensity as well. The structure factor S(h) of a grating tilted in the (x, y)-plane can be expressed in the form of Eq. (9). Sh; Dy0

sinp jh; Dy0 sinp jh; Dy0 =n

9

Further modifications have to be introduced in order to describe correctly the observed intensity decay for both signs of the scattering angle. In Eq. (10), f (Dy) represents a Gaussian distribution of Dy around the mean value Dy0, which was obtained from the parabolic fit. NL/R is a sign-dependent scaling factor, and in addition a roughness r is introduced. Z I1 h NL;R

dDy0 e

r 2 h2

f Dy Sh; Dy0

Fig. 26.6 Representative images as detected by an optical microscope during tack tests through a transparent substrate. Highly inhomogeneous structures develop on a macroscopic scale. Polymeric partition walls keep neighboring cavities separate. A polygonal, foam-like structure develops. The diameter of the punch is 2 mm.

10

431

432


Furthermore the calculated intensity was folded [Eq. (11)], B(q) being the beam profile as determined from a reference measurement. Z Ih

dh0 I1 h0 Bh

h0

11

All fits to the scattering curves are included in Fig. 26.4 b (solid lines). Except for small angles near the beam stop, the fits match well with the experimental data [39]. Therefore the assumption of a linear arrangement of voids is in good agreement with the experimental data.

26.4 Conclusion

The tack test is a powerful mechanical test to probe the mechanical response of a PSA to a given load of PSA macroscopically. The test geometry introduces a high degree of confinement which is typical for the desired type of applications. In addition it allows an analysis of the developing structural features. However, being limited by the optical resolution, these experiments have accessed the macroscopic regime only. The introduction of a new experimental technique, USAX, yields additional structural information on debonding PSA. As the scattering curves detected from the homopolymer PnBA differ qualitatively from those referring to the statistical copolymer P(EHA-co-MMA), it can be concluded that this technique has a strong sensitivity to the type of PSA investigated. The scattering signal can thus be regarded as a fingerprint of the probed material. The strongly deviating features can be due to the chemical difference between PnBA and P(EHA-coMMA), or it can be attributed to rheological properties. With respect to the elastic storage modulus G' (x), the tack test is located in a narrow frequency regime. Differences in G' (x) may lead to the observed change in the nano-object shape. A cylindrical object shape would match a more liquid-like behavior and a void object shape would match an increased elasticity. Anyway, the additional nanoscopic information will help to bridge macroscopic measurements of mechanical or rheological behavior and microscopic measurements of friction and surface forces. The scattering data support the new model of nano-cavities existing as a substructure of vertically expanded adhesive material; this shows that structure creation, wellknown at macroscopic length scales, also proceeds in the nanometer range. Scaling down the model of cavitation and swelling of cavities, the nano-bubbles or nano-cavities might be the first precursors of the larger optically probed cavities. During stretching some of the initially nano-sized cavities grow and become optically observable, whereas other nano-sized cavities remain unchanged in size. As a consequence, during debonding an apparently optically homogeneous PSA film would have nano-cavities already located inside. Whether these nano-cavities are imposed by the contact during bonding, or are due to a very local failure of the bonding during debonding, is a topic of future research.

References

This type of experiments is limited to very high-flux sources such as synchrotron radiation centers. Thus it is restricted to large-scale facilities with a limited experimental access.

Acknowledgment

We thank S. Cunis for her help during the setting-up of the BW4 beamline at the HASYLAB. T. Ittner and A. Götzendorfer helped during the USAX experiments. Additionally, we owe many thanks to R. Gehrke for his general support of the experiments at HASYLAB. We thank D. Wulff and N. Willenbacher for supplying the statistical copolymer and M. Stenert and F. Bandermann for providing the PnBA. We have obtained support from the BMBF (Förderkennzeichen 03CO333).

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27 Nondestructive Testing of Adhesive Curing in Glass–Metal Compounds by Unilateral NMR K. Kremer, B. Blümich, F.-P. Schmitz, and J. Seitzer

Abstract

NMR spectroscopy and NMR imaging require expensive equipment because the magnetic fields have to be very homogeneous and strong to measure the discrete distributions of frequencies in the parts per million range. In unilateral NMR, a magnet and the RF communication antenna are placed on the object, which can be much larger than the magnet. With permanent magnets, NMR sensors as small as a computer mouse can be built and positioned on intact objects at different places to measure T2 (the transverse NMR relaxation time), which correlates with molecular mobility. The use of the NMR-MOUSE® (a unilateral NMR scanner, the Mobile Universal Surface Explorer) for nondestructive online monitoring of adhesive curing for car windshields has been investigated. The samples studied were from heat-curing as well as from moisture-curing polyurethane (PU) rubber. The adhesives were positioned between sheets of glass (6 mm) and iron (0.8 mm). The T2 measurements took place at a depth of 7–9 mm using the CPMG (see H. Y. Carr, E. M. Purcell, Phys. Rev., 94, 630 (1954)) technique. Cured and uncured PU differed in T2 by a factor of 2. Measurements at temperatures between 10 and 45 8C showed that T2 depends on temperature. Four different rubber types could be distinguished by T2 measurements. All these measurements showed the possibility of detecting the change of molecular mobility following the curing process. For the moisture-cured and two-component PU samples, the complete curing processes were followed over time. A striking result was that full curing of moisture-curing PU takes much longer than expected. Furthermore, artificial defects in moisture-cured samples could be detected by the NMR-MOUSE® by changes in the signal amplitudes.


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27 Nondestructive Testing of Adhesive Curing in Glass–Metal Compounds by Unilateral NMR

27.1 Introduction

NMR is a standard tool for spectroscopy in chemical analysis and a noninvasive method of imaging in medicine. Usually, the object of interest is positioned inside the cavity of a magnet and irradiated with magnetic radio-frequency (RF) waves [1]. Such instrumentation is large and stationary, and the associated magnetic fields are highly homogeneous. In unilateral NMR the magnetic field is applied to the sample from one side [2]. However, small NMR devices, such as the NMR-MOUSE (a unilateral NMR scanner, the Mobile Universal Surface Explorer), can be built so that they are portable and can be carried to the object (Fig. 27.1) [3–6]. The associated magnetic fields are highly inhomogeneous and echo techniques need to be employed for measurement. Soft matter such as rubber and other polymer materials can readily be characterized by unilateral proton NMR in terms of a variety of parameters which correlate with the overall molecular mobility [7]. This way, information is obtained nondestructively about the crosslink density, the state of cure or overcure, the effect of aging, and product heterogeneity. For this reason the NMR-MOUSE can be used to optimize product development, and to monitor product and production quality. In this article, the principles of unilateral NMR are reviewed, and a concept is outlined for nondestructive quality control in rubber manufacturing.

27.2 Nuclear Magnetic Resonance (NMR) and the NMR-MOUSE

NMR is a spectroscopy technique using magnetic radio-frequency (RF) waves in magnetic fields [1]. Magnetic atomic nuclei such as the proton, 1H, are contacted in magnetic fields by the magnetic component of an electromagnetic wave. The resonance frequency of that wave is proportional to the strength of the applied field. For a B0 field strength of 0.5 T, typical for the NMR-MOUSE, the frequency for protons is nearly 20 MHz. The contact is achieved by RF exci-

Fig. 27.1 The NMR-MOUSE measuring the rubber insulation of a high-power underground cable sleeve. A defect is located at positions 1, 2, and 3 with an average of T2eff = 3.6 ms, while the regular material exhibits an average value at positions 4, 5, and 6 of T2eff = 3.8 ms. T2eff is defined in the text.

27.3 Quality Control

tation pulses a few microseconds long. In principle the decay time constant T2 of the ideal pulse response in perfectly homogeneous magnetic fields is of interest. It is called the transverse NMR relaxation time. But the inhomogeneous field of the NMR-MOUSE requires a stroboscopic observation by a train of magnetization echoes generated by many RF pulses. Although the resulting relaxation time T2eff depends on the geometry of the NMR-MOUSE and on the pulse sequence, it is in some respects similar to the ideal T2. Both relaxation times are determined by the type and the frequencies of molecular motion in the object. Typical values of T2eff are 1 s in liquids and 1 ms in rubber [3, 4].

27.3 Quality Control

A very important quantity for characterization of rubber is the crosslink density. It is proportional to the modulus at small sample deformations and can be determined on test samples under torsion in a rubber process analyzer during vulcanization. Alternatively, swelling experiments are often conducted on test samples [8]. The crosslink density can be assessed nondestructively by indentation, a principle which is used in measuring the Shore hardness. Many studies have shown that NMR is an excellent tool for characterizing rubber in detail in terms of chemical crosslink density by relaxometry, in terms of chemical structure by spectroscopy, and in terms of product heterogeneity by NMR imaging [1, 9, 10], but so far NMR has not been established as an accepted routine method in rubber testing. Nevertheless, mobile NMR of rubber with the NMR-MOUSE has the unique advantage of being nondestructive [4–7], so that it can be carried out on the final product for quality control. The advantages can be summarized as follows: · The measurement is nondestructive. · The equipment is mobile. · Information about the state of cure, formulation errors, and changes in processing parameters is obtained at all intermediate stages, from the raw material to the finished product. · The measurement is fast. · The measurement is accurate. · Execution of the measurements requires little training. · The equipment cost is competitive. Statistical averages need to be performed to obtain accurate results. This is known to apply to all kinds of measurements on rubber but is often neglected. For example, swelling investigations corroborate the inherent inhomogeneity of technical rubber and correlate well with measurements of T2eff. The same applies to mechanical testing [6]. Although theoretical models [1] exist, the correlation of T2eff with crosslink density is done in practice by calibration [1, 7], because not only the formulation but also the processing conditions and the state of cure are of importance. Moreover, it

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is possible to discriminate between incomplete cure, and curing for too long beyond reversion, in samples like natural rubber [4], because the competing reactions of chain scission and crosslinking show a crossover in the rheometer maximum and affect the chain mobility and thus the T2eff in different ways. While the rheometer torque and the degree of swelling are the same at curing times t90 and tR90, T2eff decreases through the rheometer maximum with increasing curing times. This illustrates that the information provided by NMR is similar, but in some ways also complementary, to crosslink density measurements. In unilateral NMR, a magnet and the RF communication antenna are placed on the object, which can be much larger than the magnet. With permanent magnets, NMR sensors as small as a computer mouse can be built and positioned on intact objects at different places to measure T2, which goes hand in hand with molecular mobility.

27.4 Application

The possibility of using the NMR-MOUSE® for nondestructive online monitoring of adhesive curing for car windshields has been investigated [11]. The samples studied were from heat-curing as well as moisture-curing polyurethane rubber from EFTEC AG, Switzerland. The adhesives were positioned between sheets of glass 6 mm thick and iron 0.8 mm thick (Fig. 27.2). The T2 measurements took place at a measurement depth of 7–9 mm using the CPMG technique 8 (a pulse sequence with one 908 pulse and several 1808 pulses). Fig. 27.3 shows typical CPMG measurements of an uncured and a cured sample. The parameters indicative of the material properties are extracted from the experimental data by fitting model functions such as the monoexponential function [Eq. (1)] to the measured signal, where the fit parameters are A and T2. The signal of the uncured sample decays more slowly than the signal of the cured one. Cured and uncured PUs differ in T2 by a factor of 2. at Ashort expf t=T2short g Along expf t=T2long g

1

Whether it is possible to detect different crosslink densities in different heat-curing PU adhesives was checked. The crosslinking appears at 140 8C with a 30 min

Fig. 27.2 Schematic drawing of the NMRMOUSE with a glass–PUR-metal probe.

27.4 Application

Fig. 27.3 CPMG measurements of an uncured and a cured sample.

curing time. The different crosslink densities were set by changing the isocyanate/ hydroxyl group ratio (NCO/OH). Four different ratios (NCO/OH = 0.3, 0.5, 0.7, 0.9) with different degrees of hardness were measured. As expected, the relaxation times of the uncured samples did not vary, because the molecular mobility should not have changed. After the samples had been cured, T2 was measured again. The decays were fitted by mono- and biexponential functions to find the best fit. A clear dependence of relaxation time on NCO/OH ratio was obtained. The T2 values decreased as the ratio was raised (Fig. 27.4). This showed that the molecular mobility can be detected as a function of crosslink density in the cured state. To investigate the temperature dependence of T2, it was necessary to check the heating from the RF pulses and the influence of the room temperature on the experimental setup. The temperature inside the sample increased only by 1.0 K during one experiment. A box was built to control the temperature of the setup in the range from 10 to 45 8C, which can be found in garages. Fig. 27.5 shows that T2 rises with increasing temperature. With higher temperatures, the static dipolar coupling becomes weaker so that the magnetization decays more slowly. For a temperature variation from 15 to 35 8C, T2long decreases by 5.08 ms. Hence, it is necessary to measure the temperature coefficient of T2 and to refer it to a standard temperature. A complete curing process was monitored for a moisture-curing PU adhesive. A curing time of 10–12 days is expected for this adhesive. The starting time was set to day 0. Daily measurements at the same time of day, up to day 9, showed nearly no change in T2long. Apparently, the curing process had not started in the sensitive volume at this time. The experiments showed the expected decrease in the relaxation time between days 15 and 28 (Fig. 27.5). The curing started at the sides of the sample, where the adhesive had contact with air. Inside the sample

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Fig. 27.4 Relaxation times at different formulations in the (a) uncured and (b) cured state.

the adhesive was shielded from moisture by the metal and glass sheet. Hence, the water had to ingress into the PU from the sample edges and the curing started from outside. Placing the sensitive volume of the NMR-MOUSE inside the PU sheet revealed that the curing process needed much more time than expected. When window panels are fitted, defects may occur. The detection of voids plays an important role in guaranteeing a good joint with the metal. Four samples with different air holes were prepared and measured (Fig. 27.6). Fig. 27.7 a shows the relaxation times versus the hole widths. With higher defect sizes only a small change in T2 was noticed. This change happens because the number of nuclei becomes so small in the sensitive volume that the decay is extremely modified. Hence, the fittings cannot be compared [1]. The amplitudes provided a clear effect, however. The amplitude is proportional to the amount of nuclei in the sensitive volume and hence proportional to the amount of adhesive. With an air void in the sensitive volume, the amount of PU was less and the amplitude went down.

Fig. 27.5 The complete curing process of a moisture-curing PU system.

27.4 Application

Fig. 27.6 Samples with air voids of different widths: (a) 9 mm; (b) 6 mm; (c) 4 mm; (d) 3 mm.

Fig. 27.7 Sketch of (a) the relaxation times and (b) the amplitudes versus the defect widths.

Fig. 27.7 b shows a clear dependence between the amplitude and the defect size. As the defect got bigger, the amplitude was dramatically reduced. This means that artificial defects in moisture-cured samples can be detected by observing the amplitude of the measured signal.

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Fig. 27.8 Testing the windshield of a car with the NMR-MOUSE.

27.5 Conclusion and Outlook

All measurements showed that it is possible in principle to detect by NMRMOUSE measurements the curing status of adhesives used for connecting window panels. Not only can the cured and the uncured states of the adhesive be distinguished, but different formulations can also be detected. Moreover, the curing processes can be tracked. Artificial defects are possible to detect by observing the amplitude. A striking result was that full curing of moisture-curing PU takes much longer than expected. For industrial use of this measurement system it is necessary to know the temperature coefficient of the adhesive to refer the relaxation time to a standard temperature. The next step in this project will be measurements on cars and after changing the window (Fig. 27.8). This method is applicable not only to cars: even trains or any other adhesive joints of glass to metal or of glass compounds can be tested.

Acknowledgments

The authors were supported by and worked in collaboration with EFTEC AG, Romanshorn, Switzerland.

References

References 1 B. Blümich, NMR Imaging of Materials, 2

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4

5 6

Clarendon Press, Oxford, 2000. G. A. Matzkanin, A Review of Nondestructive Characterization of Composites Using NMR, in: P. Höller, V. Hauck, G. Dobmann, C. Ruud, R. Green, Eds., Nondestructive Characterization of Materials, Springer, Berlin, 1989, pp. 655–669. G. Eidmann, R. Savelsberg, P. Blümler, B. Blümich, J. Magn. Reson. A122, 104– 109 (1996). B. Blümich, S. Anferova, K. Kremer, S. Sharma, V. Herrmann, A. Segre, Spectroscopy 18, 24–54 (2003). B. Blümich, M. Bruder, Kautsch. Gummi Kunstst. 56, 90–94 (2003). B. Blümich, S. Anferova, F. Casanova, K. Kremer, J. Perlo, S. Sharma, Kautsch. Gummi Kunstst. 57, 346–349 (2004).

7 V. Herrmann, K. Unseld, H.-B. Fuchs,

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9

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B. Blümich, Colloid Polym. Sci. 280, 738– 746 (2002). A. Kumar, R. K. Gupta, Fundamentals of Polymer Engineering, Marcel Decker, Basel, 2003. V. M. Litvinov, P. P. De (eds.), Spectroscopy of Rubbers and Rubbery Materials, Rapra Technology Ltd, Shawbury, 2002. P. Blümler, B. Blümich, Rubber Chem. Technol. (Rubber Reviews) 70, 468–518 (1997). K. Kremer, H. Kühn, B. Blümich, J. Seitzer, F. P. Schmitz, Adhäsion, 32–36 (2002).

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28 Chemical Processes During Aging in Ultra-thin Epoxy Films on Metals A. Meiser, C. Wehlack, and W. Possart

Abstract

The aging behavior of ultra-thin epoxy films (DGEBA and DETA) on metal substrates (Al, Au, Cu) has been studied in order to reveal substrate influences on chemical aging mechanisms in the interphase region that differ from bulk aging. Different environmental conditions have been applied at 40 8C for up to 100 days to separate the roles of temperature, H2O, and O2 in aging. The film thickness was varied (60–650 nm) in order to identify the influence of the metal substrate on the aging behavior in the interphase. Quantitative evaluation of the IR spectra provided the following results. Contrary to the bulk, the initial epoxy group conversion of the thin films was incomplete. During aging, post-curing took place that was catalytically enhanced by metal substrates in a specific way and by the presence of water, and characteristic aging bands appeared. Their interrelation was more complex than described in the literature. It was shown that the metallic substrates, the aging conditions, and the epoxy film thickness influenced the aging effects and their kinetics. With these experimental results, chemical aging mechanisms are discussed.

28.1 Introduction

The aging of adhesives and lacquers affects the durability of structural bonds and coatings. It depends on environmental influences such as humidity, temperature, irradiation, the surrounding media, and the network state of the polymer. In addition, the network properties and therefore the aging behavior are influenced by substrates that provoke the formation of an interphase during network formation. The properties of the interphase differ from the bulk and it is often the weak spot of structural adhesive joints. Hence, to improve the lifetime of such a joint it is crucial to understand the chemical processes both in the bulk polymer and in the adhesive–adherend interphase under given enviAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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28 Chemical Processes During Aging in Ultra-thin Epoxy Films on Metals

ronmental conditions. In this paper, the chemical aging behavior of aminecured epoxy networks is examined because of their high technical relevance as structural adhesives and protective coatings. The aging of bulk epoxy networks has been studied by various techniques such as FTIR spectroscopy [1–4], gravimetric analysis [1, 3, 5, 6], mass spectroscopy [1], DSC [2], or XPS [6]. The aging conditions applied vary broadly from thermo-oxidation at elevated temperature (100–250 8C) [1–3, 5, 6] or photo-oxidation [1, 2] to humid or environmental conditions [4, 7]. FTIR investigations [1–4] show the rise of new IR bands at 1660–1670 cm–1 and 1720–1730 cm–1 in an oxidized region near to the sample surface. The bands are explained by amide and carbonyl formation due to radical oxidation mechanisms initiated by the elevated temperature or UV irradiation [1, 2, 4]. Also, backbone cleavage is held responsible for the observed decrease in glass transition temperature [2]. However, direct investigation of the interphase properties inside an adhesive bond is very difficult, for experimental reasons. In contrast, ultra-thin polymer films on metal substrates provide access for a number of analytical techniques. Hence, to some extent, they are suited to monitoring of interphase properties and processes. By varying the film thickness, it is possible to discriminate the influence of the metal substrate on the aging behavior in the interphase. Only a few workers have investigated the influence of metal substrates on network formation and on aging for thermosetting films (30–1500 lm) on metal substrates (Al, Ti, Au [8–11]; Al, Cu, steel [12]) or in bulk samples made of metal-modified monomers [8–11] or filled with metal powders [13, 14]. These studies reveal the formation of a chemical interphase with properties quite different from the bulk. A decreased glass transition temperature and reduced amine conversion are observed. These findings are explained by a preferential adsorption of the hardener and by a partial dissolution of the metal hydroxide layer due to formation of amine–metal complexes [8–11]. Thermo-oxidation of amidoamine-cured epoxy films on metals and on copper oxide-filled brominated epoxy resins at elevated temperatures (150–165 8C) leads to similar IR bands to those observed for bulk oxidation [12] but degradation is accelerated on the metal substrates, especially by copper [12, 13]. However, the aging behavior of thin epoxy films at moderate temperature and under the additional influence of humidity is unknown, and the aging behavior of ultra-thin films that are best suited to monitoring of substrate influences has not been examined so far. This paper reports the study of the aging behavior of ultra-thin epoxy films on PVD layers of Al, Au, and Cu under varying environmental conditions by FTIR external reflection absorption spectroscopy (ERAS), to reveal influences of interphase interactions on crosslinking and on chemical aging processes at moderate temperature.

28.2 Experimental

28.2 Experimental 28.2.1 Sample Preparation

The epoxy system examined consisted of the diglycidyl ether of bisphenol A (DGEBA, Dow Plastics D.E.R. 332) and the curing agent diethylenetriamine (DETA, Dow Plastics D.E.H.). The chosen mass ratio, DGEBA/DETA = 100 : 14, corresponds to a slight excess of amine hydrogen compared with epoxy rings. Both components were mixed and stirred in the molten state at 55 8C for 5 min in a closed glass vessel. A subsequent pre-curing step of 1 h at room temperature prevented a quantitative loss of the volatile hardener during and after thin-film preparation. The metal substrates were obtained by physical vapor deposition of ca. 100 nm of pure Al, Cu, or Au onto Si (100) wafers. The resulting polycrystalline metal surfaces were covered by their natural oxides/hydroxides (except for Au) and an organic adsorption layer. Ultra-thin epoxy films with thickness dEP of 60, 125, 250, and 650 nm were spincoated in CO2-reduced, dried air (dew point –70 8C) from the appropriately diluted prepolymer solution in methyl ethyl ketone (MEK). After film preparation, the specimens were stored in the same atmosphere at about 23 8C for 72 h. This room temperature (RT) curing was followed by a “post-curing” step at 120 8C in argon for 1 h. In order to check reproducibility, two to four samples with the same substrate and film thickness were prepared separately for each aging regime. 28.2.2 Aging Conditions

In the dark and at a moderate aging temperature of 40 8C, different environmental conditions were applied to the thin-film samples to discriminate between the roles of temperature, H2O, and O2 in aging: purely thermal effects in a dry argon atmosphere, hygro-thermal aging in humid Ar (90% r.h.), thermooxidative aging in dry air, and hygro-thermo-oxidative aging in humid air (90% r.h.). After 100 days under constant aging conditions, some of the dry-aged samples were humidified (90% r.h.), whereas hygro-thermally aged samples were re-dried at room temperature in order to reveal reversible processes. During aging, chemical modifications were studied by FTIR external reflection absorption spectroscopy (IR-ERAS) with p-polarized light at a 708 angle of incidence in the mid-IR range (4000–400 cm–1). Before measurement, the samples were conditioned in the laboratory atmosphere for 30 min to equilibrate their water content. Depending on film thickness, 200–500 scans were co-added to attain a reasonable signal-to-noise ratio. The film spectra were divided by the spectra of the bare metal substrate (reference) to ratio out the influence of the spectrometer function (e.g., light source, gas atmosphere in the light path).

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28.3 Results and Discussion 28.3.1 Crosslinking

For the given conditions, curing in this epoxy system proceeds by crosslinking of oxirane with primary and secondary amine groups which results in hydroxyl groups. A quantitative analysis of amine group consumption and of hydroxyl group formation is difficult due to the fact that their IR bands overlap or are hidden by other strong bands in the mid-IR range. Fortunately, the consumption of oxirane groups can be followed by the oxirane stretching band at 915 cm–1. In order to compare this consumption quantitatively as a function of IR time and film thickness, the spectroscopic degree of oxirane conversion UEP is calculated from the band intensity I915 (peak height) of the band at 915 cm–1, which is normalized to the intensity I1510 (peak height) of the phenylene ring stretching vibration at 1510 cm–1 [Eq. (1)]. IR UEP tcure 1

I915;norm tcure I915;norm tcure 0

with

I915;norm tcure

I915 tcure I1510 tcure

1

In general, the comparison of spectra of films with different thickness is problematic due to optical influences on the resulting spectra [15]. Fortunately, simulaIR tion of thin-film spectra with constant bulk-like properties proves that UEP values vary by less than 3% for the thickness range studied, due to the changed optical situation [16]. Therefore, a direct quantitative comparison is possible within the limits of this small systematic error. In addition, the comparison between simulated bulk-like spectra and measured thin-film spectra is suited to revealing characteristic influences of the substrate and the interphase. Band intensities still depend not only on the concentration but also on the extinction coefficient of the eigenvibration considered. The extinction coefficient is difficult to determine and it varies with the chemical environment. As a result, IR UEP is a quantitative measure for the chemical conversion, but the two quantities are not identical. IR Before aging, the initial UEP values are a function of epoxy thickness (see IR Fig. 28.1): the thinner the film, the lower the UEP , although the curing temperature and time are more than sufficient for complete epoxy consumption in the bulk. Hence, the epoxy network remains incomplete in the post-cured films. Supported by other experimental results, it is concluded that the metal substrates induce the formation of a heterogeneous interphase with a locally varying reaction rate during network formation [17]. As this heterogeneity is more pronounced in the vicinity of the interface, the overall degree of conversion is weaker for thinner films due to the spatial separation of the reaction partners. The residual reactive groups might undergo further crosslinking or competing reactions during aging and these effects should be stronger in thinner films.

UIR EP [%]


Fig. 28.1 Initial spectroscopic degree of oxirane conversion UIR EP for epoxy films on Au, Al, and Cu prepared for hygro-thermo-oxidative aging (preparation 1) and for thermal aging (preparation 2).

Even the 650 nm layers do not reach the epoxy conversion in the bulk, which is almost 100% after post-curing [18]. Hence, this substrate-induced inhibition of IR the curing reaction extends over several hundred nanometers. On Cu, UEP is IR slightly increased in comparison with Al and Au, which show comparable UEP curves. IR Note that the UEP curves shift for the two sets of samples prepared on two different days. This indicates the limits of sample reproducibility due to ambient temperature fluctuations that affected the kinetics and final epoxy conversion. Therefore, different preparation series, and hence aging regimes, may be compared only qualitatively. In future work, curing conditions will have to be controlled even more precisely. Fig. 28.2 depicts the development of the oxirane concentration with aging time tage in dry Ar and moist air for 125 nm and 650 nm epoxy films as an example. IR In dry Ar (Fig. 28.2 a), UEP tage is similar on Al and Au but different on Cu. Under dry conditions on Al and Au, the residual oxirane is consumed completely within 49 days in thin layers, while only a small change is observed in the 650 nm films over the full 100 days of the experiment. This influence of dEP could indicate a smaller size of heterogeneities and an increased polymer mobility near the substrate that would facilitate diffusion during aging. On Cu in dry Ar, the kinetics of oxirane consumption tend to be slower in all films and the fiIR nal UEP tage falls behind the values on Al and Au. These findings can be related to Cu2+ ions detected by XPS and TOF-SIMS in these films [19]. Notable

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UIR EP [%]


UIR EP [%]

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Fig. 28.2 Spectroscopic degree of oxirane conversion UIR EP for epoxy films (125 nm and 650 nm) on Au, Al, and Cu under (a) thermal (Ar, 40 8C) and (b) hygrothermo-oxidative aging (air, 40 8C, 90% r.h.).


Cu–DETA complexation in the epoxy, especially in the interphase, reduces the concentration of reactive amine hydrogen and this causes slower kinetics and a smaller degree of conversion. In a humid environment (Fig. 28.2b), all the epoxy films pick up water, which plasticizes the network and acts catalytically on the oxirane conversion. Hence, IR UEP tage grows faster than under dry conditions and it depends less on film thickness. All the films on all the substrates reach almost full oxirane converIR sion within 100 days. Films on Au and Al show the same kinetics and final UEP values for a given thickness, while on Cu the kinetics of oxirane conversion tend to be slower again. There could be several reasons for the oxirane conversion observed during aging: further crosslinking with residual amine groups with formation of hydroxyl groups, homopolymerization with the present hydroxyl groups with formation of ether groups, or a ring-opening reaction with water molecules without any byproduct. Each reaction would leave its fingerprints in the IR spectra, especially in the regions of hydroxyl and ether bands. In particular, comparison of oxirane consumption and hydroxyl band development is helpful for the identification of the oxirane conversion reaction. Therefore, the hydroxyl band intensity at 3500 cm–1 (peak area) is normalized to the intensity I1510 (peak height) of the phenylene ring stretching vibration at 1510 cm–1 [Eq. (2)]. IOH;norm tcure

IOH tcure I1510 tcure

2

IR These data are compared with UEP in Fig. 28.3 for the Au substrate as an examIR ple. For all aging conditions and all samples, IOH,norm and UEP develop quite synchronously with tage. Further, no ether band formation (asymmetric stretching at ca. 1140 cm–1) is detected during aging. Therefore, homopolymerization is not the cause of the oxirane consumption during aging. It is dominated by the ongoing reaction of oxirane with amine groups during hydroxyl group formation. This conclusion is supported by the fact that the moderate aging temperature is far too low for the formation of ether crosslinks [20].

28.3.2 Additional Aging Effects

According to Fig. 28.4, two new bands at 1738 cm–1 and 1660 cm–1 are present just after post-curing, especially in thinner films and on the Cu substrate. No such band formation is observed for bulk samples. This result indicates that post-curing itself leads to remarkable chemical modifications in thin films on metals. Since these effects depend on dEP and on the substrate material, they should originate from catalytic enhancement of specific reactions that compete with the network formation. The solubility of the Cu2+ ions that are most likely to catalyze competing reactions explains why pro-

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Fig. 28.3 Spectroscopic degree of oxirane conversion UIR EP and hydroxyl band intensity IOH,norm for (a) 125 nm and (b) 650 nm epoxy films on Au during thermal or hygrothermo-oxidative aging.


Fig. 28.4 “Aging region” in IR-ERA spectra of (a) 650 nm and (b) 125 nm epoxy films on Au, Al, and Cu with new bands after post-curing.

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nounced bands at 1738 cm–1 and 1660 cm–1 appear on the Cu substrate even in 650 nm films, whereas no intensity is found for 650 nm films on Al and Au. Depending on the aging regime, the IR bands at 1738 cm–1 and 1660 cm–1 change and new bands appear at 1725 cm–1 and 1600 cm–1 on all three metal substrates. The exact position, width, and temporal development of these bands depend on dEP, on tage, and on the metal. A quantitative evaluation of their intensities proves to be difficult as the bands overlap with phenylene bands. Nevertheless, qualitative consideration reveals some interesting facts about the complex picture (see Fig. 28.5 for a selection of data). Under the different aging regimes, the bands do not develop synchronously. Hence, they indicate more than one chemical reaction with more than one product. In general, the bands themselves and the differences in intensity due to the substrate influence become more significant with decreasing dEP (Fig. 28.5 a–d). This can have two reasons. First, specific catalytic effects of the metal substrates and ions affect the aging reactions close to the metal substrate (thin films) more strongly than in thicker films. Second, thinner films possess a higher residual content of oxirane and amine than 650 nm films. The higher the concentration of these reactive species, the more they can be involved in chemical aging reactions that compete with crosslinking. As the reduced initial epoxy conversion is also attributed to a substrate influence, both reasons (catalytic cure and competing reactions) for the stronger IR bands in thinner films can be traced back to substrate effects. Films on Au and on Al show approximately the same IR band evolution, whereas on Cu it is in part more pronounced, in part weaker (Fig. 28.5). Again, this specific behavior of films on Cu substrates could be due to a catalytic or an inhibiting effect of diffused Cu2+ ions on the various aging mechanisms. Additionally, copper oxides can act as an effective oxidation–reduction couple, thus accelerating oxidation reactions [12]. However, the environmental conditions are the crucial factor for chemical modifications during aging. Water has great impact on the IR spectra whereas the oxygen content appears to be less important. Therefore, only spectra of samples aged in dry Ar and moist air are shown as examples for dry and moist aging conditions in Fig. 28.5. Under dry aging conditions (Ar or air), a new chemical species with an IR band at 1725 cm–1 is the most remarkable aging product on all the substrates (Fig. 28.5 e, f; results for dry air not shown). Its rate of formation and final concentration are higher in thinner films. Hence, its production must be either catalyzed by the metal substrates or supported by the minor initial oxirane conversion. Compared with Al (and Au), its formation is slowed down on Cu. In the course of aging in dry Ar, the band height goes through a maximum in thinner films on Al (and Au) indicating a competing degradation process for the product due to its instability or reactivity. In addition to the pronounced band at 1725 cm–1, a broad spectral background under the phenylene bands grows slightly but steadily under dry conditions. Difference spectra reveal that this background is due to a broad band at 1700–1550 cm–1 with a maximum at 1600 cm–1.


Fig. 28.5 “Aging region” in IR-ERA spectra of thin epoxy films (125 nm, 650 nm) on Al and Cu during thermal or hygro-thermo-oxidative aging and subsequent humidification or re-drying.

In humid environments (air or Ar), aging leads to a pronounced formation of products with bands at 1660 cm–1 and 1600 cm–1 (Fig. 28.5 a–d). Obviously, their formation is stimulated by the presence of water. The steady increase of the 1660 cm–1 band clearly depends on the metal substrate and is most pronounced on Al and Cu, whereas only slight changes are found on Au. Thus, specific sub-

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Fig. 28.5 c and d

strate properties such as the acidity of the hydroxide layers on Al and Cu stimulate its production. The second product (1600 cm–1) is formed without a significant substrate dependence. It is degraded in humid air on Al and Au. No band at 1725 cm–1 is visible in moist aged films. Instead, the band at 1738 cm–1 that appeared after post-curing increases weakly in moist air in the initial stages of aging. This result indicates that the formation of this species is favored by higher temperatures and by the presence of water. At a later stage, the band inten-


Fig. 28.5e and f

sity at 1738 cm–1 decreases on some samples, i.e., the product can be degraded under humid conditions. Re-drying of moist aged films has no significant impact on their chemical state (Fig. 28.5 c , d). In contrast, humidification of thermally aged samples leads to drastic changes in the IR spectra (Fig. 28.5 e, f). The band at 1600 cm–1 grows, but only on Al and Au. The pronounced band at 1725 cm–1 disappears so that the weaker initial band at 1738 cm–1 becomes visible, as is the case for moist

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aged samples. Hence, the aging product (1725 cm–1) is almost fully converted under the influence of water. The band at 1660 cm–1 grows simultaneously. The coincidence of the two phenomena suggests a sequence of reactions: first, an intermediate (at 1725 cm–1) is formed which is quickly converted to the final product (1660 cm–1) as soon as water is present. In dry conditions, the intermediate is formed but not degraded because of the absence of water. 28.3.3 Band Assignment and Chemical Aging Processes

In summary, bands at 1738, 1725, 1660, and 1600 cm–1 appear in the IR spectrum during aging. Their evolution depends on the environmental conditions and on the nature of the metal surface. As the next step, the IR bands have to be assigned to new chemical groups in order to deduce the chemical aging mechanisms in the epoxy films (see Table 28.1). This is difficult, however, because band assignment is not unambiguous and different aging mechanisms are possible. The hydroxyl groups formed during crosslinking can be oxidized to carbonyl groups, especially at the elevated temperature during sample preparation (see Scheme 28.1). This process can cause the band at 1738 cm–1 that appears after post-curing. In addition, the oxidation of carbonyl groups is catalyzed by copper in common chemical practice and this would explain why the band is most pronounced on this substrate. In the literature on aging in the epoxy bulk, IR bands at 1730 and 1660 cm–1 are attributed to carbonyl and amide groups that result from radical oxidation of

Table 28.1 Conditions of formation, stability, and band assignment for the aging bands observed in thin epoxy films on Al, Au, or Cu.

IR band [cm–1]

Formation conditions

Stability

Band assignment

1738

during cure on Cu, weaker degradable on Al, Au, and Cu carbonyl on Al and Au; further growth during aging in humid during aging in humid air atmosphere

1725

during aging in dry atmosphere on Al and Au, slow and weak on Cu

converted with H2O in humid atmosphere

carbonyl

1660

during cure on Cu, Al, and Au (weak); further growth in humid atmospheres

stable

amide/oxime

1600

during aging in dry and stronger in humid atmospheres on Al, Cu, and Au

stable on Cu, degradable on Al and Au during aging in moist air

amino salt


R1

R1

CH2 CH2

Scheme 28.1 Oxidation of hydroxyl groups formed during crosslinking.

backbone C atoms initiated by elevated temperatures (about 100–250 8C) or by UV irradiation at moderate temperature (see, e.g., Refs. [1–6, 12, 13]) as depicted in Scheme 28.2. Alkyl radicals and molecular oxygen form peroxy radicals. They are converted into hydroperoxides by proton transfer from some other proton donating group in the network. The resulting hydroperoxides tend to dissociate. The final oxidation products, such as carbonyls (ketones or aldehydes) and amides, are formed by peroxy decomposition (with or without network cleavage). Due to the low chemical selectivity of peroxy radicals, a variety of further products in different chemical environments could be created. Thin-film aging could be based on such radical oxidation processes of the polymer backbone because the products would explain the observed IR bands around 1730 cm–1 and at 1660 cm–1. Nevertheless, some particularities of the actual experiments and of thin-film chemistry have to be taken into account. On one hand, initiation by light irradiation is excluded by the dark aging conditions, and the moderate aging temperature of 40 8C is far below common initiation temperatures in thermo-oxidative degradation experiments described in the literature. Therefore, thermal radical formation and oxidation kinetics are at least slowed down considerably. On the other hand, the metal ions present in the interphase can catalyze the dissociation of hydroperoxides, leading to the formation of new peroxy radicals. The stronger band intensities at 1738 and 1660 cm–1 during aging in films on copper and in part on aluminum compared with gold would support such metal-specific catalytic effects.

OH H2C CH2

Scheme 28.2 Oxidation products of a typical epoxy backbone segment as suggested in Refs. [1–6, 12, 13].

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However, the radical oxidation of backbone C atoms described does not explain the observed influence of incomplete epoxy conversion on the aging behavior of epoxy films. Contrary to the bulk or to thick-film samples examined in the literature, we found crosslinking to be far from complete in thin films before aging. While the residual epoxy groups are insusceptible to oxidation or other side reactions under the chosen aging conditions, the residual amine groups can undergo additional reactions that compete with the curing reaction. Primary and secondary amine groups are highly reactive and easier to oxidize than backbone methylene groups. The correlation between the reduced initial oxirane conversion in thinner films and the more pronounced aging effects suggests that oxidation of the remaining reactive amine groups plays a key role in aging of thin films. In chemical practice, primary amine groups are easily oxidized to oximes [21]. The most convenient oxidant cited is hydrogen peroxide at room temperature, but catalytic oxidation with molecular oxygen is also reported (Scheme 28.3). In thin films, the metal substrate could act as such a catalyst. Because oximes absorb light at 1670–1649 cm–1 [22, 23], they can be the cause of the aging band at 1660 cm–1. However, this reaction path does not fully explain why this band is observed only under humid aging conditions. Secondary amines are less susceptible to oxidation. Nevertheless, with an excess of oxidants and under catalytic conditions, oxidative backbone scissions lead to the formation of aldehyde and oxime species [21] as depicted in Scheme 28.4. They absorb in the IR range of the aging bands observed at 1730 cm–1 and at 1660 cm–1. The metal substrates could act as catalysts again. Indeed, no synchronous development of band intensities is observed. Therefore, that mechanism should only complement the other reactions. Tertiary amines are even more difficult to oxidize and temperatures of more than 100 8C are needed to decompose the resulting amine oxides into products that absorb in the 1750–1550 cm–1 range via Cope elimination [24, 25]. Thus the oxidation of tertiary amines is hardly responsible for the observed aging phenomena. Bands at 1600 cm–1 are attributed to the formation of primary amines by network scission at high temperature [12]. In our case, the ongoing cure of residual reactive amines should rather lead to a decrease in primary amine concentration. Therefore, the band must have another origin, such as the protonation of primary and secondary amines to form amino salts or the formation of carboxyl

Scheme 28.3 Oxidation of primary amines to oximes.


Scheme 28.4 Oxidation of secondary amines to aldehydes and oximes according to Ref. [21].

acids and salts. These species absorb at 1600 cm–1 [23]. Nevertheless, neither the band assignment nor a reaction mechanism can be proposed unambiguously for this band so far. Summing up, the observed growth of bands during aging is probably due to the formation of carbonyl groups in a different chemical environment (1738 and 1725 cm–1), amide or oxime groups (1660 cm–1), and amino salts or carboxyl species (1600 cm–1). Indeed, neither radical backbone oxidation nor oxidation and conversion of amine groups can fully explain the independent development of the bands around 1730, 1660, and 1600 cm–1. Nevertheless, the correlation between the residual amine content and the strength of the aging effects favors the hypothesis that amine oxidation and conversion form a major reaction path. However, things are made more complex by the degradation observed for most of these products. The decrease in the prominent band at 1725 cm–1 during thermal aging and subsequent humidification could be explained by a reaction of carbonyl groups produced, with residual amine groups. In fact, the synchronous rise of the IR band at 1660 cm–1 during humidification could be explained by the presence of the resulting imine and enamine groups (see Scheme 28.5).

Scheme 28.5 Conversion of carbonyl and amine groups into imine or enamine groups.

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This reaction might also be responsible for the slight decrease in the band at 1738 cm–1 in a humid environment. However, the cause of the observed influence of water on this conversion remains unclear since its presence should push the equilibrium back to the educts instead of promoting the carbonyl conversion. Finally, the effects of the substrate surfaces as possible oxidants or catalysts and the role of adsorbed water or oxygen are even more complex and cannot be understood from FTIR spectra alone.

28.4 Conclusion

The aging behavior of ultra-thin epoxy films on metal substrates at moderate aging temperatures has been investigated for the first time by FTIR ERA spectroscopy. The results reveal that the crosslinking and aging differ significantly from those in the bulk. First of all, the metal substrates inhibit full crosslinking during film preparation in conditions that provide complete chemical conversion in the bulk. The corresponding chemical interphase extends over several hundred nanometers. Crosslinking of epoxy and amine groups goes on during aging and it is accelerated in thin films and by water. Furthermore, a complex interrelationship between multiple chemical aging mechanisms leads to aging products with IR bands at 1738, 1725, 1660, and 1600 cm–1. Their formation kinetics, final concentration, and chemical stability are influenced significantly by the metal substrate, by the initial content of residual oxirane/amine groups, and by the aging conditions, especially by the presence of water. According to the spectroscopic data we suppose that reactions of residual amine groups form the major path of chemical modification in epoxy films on metals during aging. It is likely that primary amine groups are oxidized by O2 to oxime groups. The secondary amines are expected to be oxidized to aldehyde and oxime groups. This process causes damage to epoxy network chains and should be relevant to the mechanical properties. Finally, the carbonyls created could also react with remaining primary and secondary amine groups, thus forming imine and enamine species. Indeed, the crucial role of water is not understood so far, and further mechanisms based on radical backbone oxidation cannot be excluded. On Cu substrates, the aging behavior of the epoxy films is different in quantitative terms from the situation on Al and Au. This is related to catalytic or inhibiting effects of dissolved Cu2+ ions. Our results show that thin-film samples are suited to monitoring specific substrate influences on the aging behavior in the interphase. Nevertheless, further experiments are needed to correlate the observed evolution of IR bands definitely with chemical reaction mechanisms.

References

Acknowledgments

The financial support of the German Federal Ministry for Education and Research (BMBF) is gratefully acknowledged.

References 1 Damian C., Espuche E., Escoubes M.,

2 3 4 5

6 7

8

9

10 11

12 13 14 15

Polym. Degrad. Stabil. 2001, 72(3), 447– 458. Bellenger V., Verdu J., J. Appl. Polym. Sci. 1985, 30(1), 363–374. Colin X., Marais C., Verdu J., Polym. Testing 2001, 20(7), 795–803. Luoma G. A., Rowland R. D., J. Appl. Polym. Sci. 1986, 32, 5777–5790. Barral L., Cano J., López J., López-Bueno I., Nogueira P., Abad M. J., Ramírez C., Europ. Polym. J. 2000, 36(6), 1231–1240. Buch X., Shanahan M. E. R., J. Appl. Polym. Sci. 2000, 76, 987–992. Fata D., Alterung eines Epoxidklebstoffes unter dem Einfluss von Wasser und erhöhter Temperatur, Diplomarbeit, Universität des Saarlandes, 2001. Bentadjine S., Petiaud R., Roche A. A., Massardier V., Polymer 2001, 42(14), 6271–6282. Bentadjine S., Roche A. A., Bouchet J., Epoxy–diamine adhesives on metals: the interphase formation and characterization, in: Adhesion Aspects of Thin Films (Ed.: Mittal K. L.), Vol. 1, 2001, pp. 239– 260. Bouchet J., Roche A. A., J. Adhes. 2002, 78, 799–830. Roche A. A., Bouchet J., Bentadjine S., Int. J. Adhes. Adhesives 2002, 22(6), 431– 441. Hong S. G., Polym. Degrad. Stabil. 1995, 48, 211–218. Hong S. G., Wang T. C., J. Appl. Polym. Sci. 1994, 52, 1339–1351. Hong S. G., Wang T. C., Thermochim. Acta 1994, 237, 305–316. Wehlack C., Möglichkeiten von IR-Spektroskopie und Spektrensimulation für die Untersuchung dünner Epoxidklebstoffschichten auf Aluminium-Substraten, Diplomarbeit, Universität des Saarlandes, 2001.

16 Petersen C., Lichtmikroskopische und

17

18

19

20 21

22

23 24

25

IR-spektroskopische Untersuchung von dünnen Epoxidschichten auf Metallen, Diplomarbeit, Universität des Saarlandes, 2002. Possart W., Krüger J. K., Wehlack C., Müller U., Petersen C., Bactavatchalou R., Meiser A., Formation and Structure of Epoxy Network Interphases at the Contact to Native Metal Surfaces, C.R. Chimie 2005, in press. Meiser A., Alterung eines Epoxidklebstoffes im Bulk und in dünnen Schichten auf metallischen Oberflächen, Diplomarbeit, Universität des Saarlandes, 2003. Dieckhoff S., Wilken R., Noeske M., Adhäsions- und Alterungsmechanismen in Polymer–Metall-Übergängen, BMBF-Project No. 03D0074 Report, Bremen (Ed.: Dieckhoff, S.), 2004. Wang X., Gillham J. K., J. Appl. Polym. Sci. 1991, 43, 2267–2277. Metzger H., Oxime, in: Methoden der Organischen Chemie (Houben-Weyl), Band XI/1: Stickstoffverbindungen II (Ed.: Müller E.), 4th edn., Georg Thieme, Stuttgart, 1957. Lin-Vien D., Colthup N. B., Fateley W. G., Grasselli J. G., Handbook of Infrared and Raman Characteristic Frequencies of Organic Molecules, Academic Press, San Diego, 1991. Socrates G., Infrared Characteristic Group Frequencies, John Wiley, Chichester, 1994. Latscha H. P., Kazmaier U., Klein H. A., Organische Chemie – Chemie-Basiswissen II, Springer, Berlin, 2002. Freytag H., Möller F., Pieper G., Söll H., Umwandlung von Aminen, in: Methoden der Organischen Chemie (Houben-Weyl), Band XI/2: Stickstoffverbindungen II (Ed.: Müller E.), 4th edn., Georg Thieme, Stuttgart, 1957.

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29 Depth-Resolved Analysis of the Aging Behavior of Epoxy Thin Films by Positron Spectroscopy J. Kanzow, F. Faupel, W. Egger, P. Sperr, G. Kögel, C. Wehlack, A. Meiser, and W. Possart

Abstract

During recent decades positron annihilation spectroscopy has become a very powerful tool for the investigation of polymers. In particular, positron annihilation lifetime spectroscopy (PALS) yields valuable information about free volume and related properties. Moreover, special chemical information can be obtained. Now advances in positron beam technology also allow investigations of thin polymer films and surface regions. In this paper, we report the use, for the first time, of PALS to elucidate aging mechanisms in thin epoxy films, based on depth-resolved investigations of the epoxy films exposed to two different aging conditions. We also consider the results of IR external reflection absorption spectroscopy (IR-ERAS) and X-ray photoelectron spectroscopy (XPS) depth-profiling of the elemental composition. This additional information enables us to clarify structural modifications due to aging. Nitrogen depletion and a decrease in free volume were observed, especially in the near-surface region of the thin epoxy films.

29.1 Introduction

The durability of the properties of epoxy adhesives is often limited by aging effects. Depending on the aging conditions, such as humidity and temperature, various structural properties of the adhesives can be strongly influenced. However, the underlying mechanisms in the early stage of aging have not been understood sufficiently. Positron annihilation lifetime spectroscopy (PALS) allows the quantitative investigation of the polymer free volume [1, 2]. Additionally, the PALS beam technique makes a direct depth resolution possible, by implanting the probe – the positron – within a definite sample depth interval depending on the positron kinetic energy [3]. It is one of the very few nondestructive techniques for investiAdhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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29 Depth-Resolved Analysis of the Aging Behavior of Epoxy Thin Films by Positron Spectroscopy

gating structural properties of samples with a depth resolution on a nanometer scale. Its additional high sensitivity to electron acceptor groups provides access to the very early stage of aging and fatigue. In this paper, 650 nm thin epoxy films (diglycidyl ether of bisphenol A (DGEBA) and diethylenetriamine (DETA)) are prepared on metal substrates (Al, Au, Cu PVD layers) and cured for 48 h at room temperature followed by 1 h of post-curing at 120 8C. Two different aging conditions are applied at 40 8C for 100 days: hygro-thermal aging in air with a relative humidity of 90% and pure thermal aging in a dry argon atmosphere. Depth-resolved PALS investigations of the polymer films prior to and after aging are presented. The results are compared and analyzed with respect to depth-resolved X-ray photoelectron spectroscopy (XPS) analysis of the chemical composition as well as in-situ studies with IR external reflection absorption spectroscopy (IR-ERAS) during aging.


PALS is a well-established technique for studying the free volume of polymers [2] by analyzing the lifetime so-Ps of ortho-positronium, a metastable hydrogenlike state consisting of a positron and an electron: Ortho-positronium is preferably located in free-volume voids due to the low potential energy in these voids. There, so-Ps is generally limited in polymers by a pick-off annihilation with weakly bonded electrons at the walls of the free-volume voids. Therefore, the average lifetime so-Ps increases with increasing free-volume void sizes. Furthermore, average free-volume void sizes can be extracted from so-Ps by using a semi-empirical equation calibrated with polymers and zeolites of known freevolume void sizes [1]. This equation can be used for most of the polymers with void sizes in the subnanometer range. The positronium formation probability Io-Ps is strongly reduced by the presence of electron acceptor groups in the polymer, since these groups trap free electrons that have been excited in the spur of the injected positrons, which otherwise could have formed positronium together with the positron [2, 4, 5]. The positron lifetime spectra are decomposed into three discrete lifetime components using the Patfit88 package [6]. While two short-lifetime components reflect para-positronium and free positron annihilation, the longest lifetime represents so-Ps. The behavior of so-Ps can be used to identify structural polymer properties such as glass transition [7] or degree of crystallinity [8], or the curing process of thermosets [9, 10]. The calculation of the free-volume void radius rvoid has also been used successfully in the quantitative analysis of gas diffusion in polymers [11, 12]. The very sensitive parameter Io-Ps has been applied to characterize, for example, early stages of the decomposition of polymer blends [13] and the behavior of metal–polymer mixtures during ball-milling [14]. The positron beam technique at the University of the Bundeswehr, Munich, enables depth-resolved investigation of polymers by implanting the posi-

29.3 Results

tron within a definite sample depth interval. The average implantation depth d depends nonlinearly (d * E1.71) on the positron energy E. The depth resolution is approximately half of the average implantation depth [3]. Infrared spectroscopy experiments are performed at the University of the Saarland, Saarbrücken, using the external reflection absorption spectroscopy (ERAS) technique. Absorption peaks are attributed to molecular eigenvibrations which in some cases are localized at specific atom groups within the molecules. Band intensities relate to the concentration of the respective atom groups. Details of the quantitative analysis can be extracted from Refs. [15] and [16]. Depth-resolved XPS analysis of the chemical composition uses an Ar ion beam microsectioning technique with a high depth resolution of the order of 2 nm. Ions with 160 eV beam energy are used to minimize sputtering artifacts arising from a high-energy transfer [17, 18]. Charge effects are eliminated by means of a neutralizer filament. After defined time intervals, the sputtering process is interrupted to measure the XPS intensity of the nitrogen N 1s as well as carbon C 1s and oxygen O 1s spectra of the polymer surface freshly created by sputtering. The sample stays in high vacuum during the whole depth profiling process. The XPS investigations are performed using a hemispherical electron analyzer and a non-monochromatized Al/Mg Ka X-ray source with photoelectron collection normal to the sample surface. The XPS peak areas are obtained after a Shirley-type background removal.

29.3 Results

The results of the PALS investigations of the aging behavior of thin epoxy films are presented here. The decisive parameters of the PALS analysis for polymers are the formation probability Io-Ps of ortho-positronium and its average lifetime so-Ps. While Io-Ps predominantly represents the effectivity of the electron acceptor groups in the polymer, so-Ps is directly correlated with the sizes of the freevolume voids. First, an example of a PALS analysis of an unaged epoxy film will serve as a reference for the investigations on aging. The dependence of Io-Ps and so-Ps on the average positron implantation depth (equal to the average depth in the sample) is presented, then the effects of the aging conditions mentioned above on Io-Ps are shown by the example of thin epoxy films on an Au substrate. The effects on so-Ps are presented in the same way. Subsequently, we show the differences in the aging behavior of the epoxy films on different metal substrates, and finally, we present additional investigations on the epoxy films with complementary experimental methods, which enable us to give an explanation of the PALS results.

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29.3.1 PALS Investigation of an Unaged Epoxy Film

Fig. 29.1 shows the PALS analysis of a 650 nm epoxy film on an Au substrate, which demonstrates that the dependence of Io-Ps on the average positron implantation depth. Io-Ps is quite constant at 28% over a depth interval between 50 and 400 nm. Above 400 nm, Io-Ps is significantly reduced. Some of the positrons with implantation depths above 650 nm penetrate into the metal substrate or even deeper into the silicon wafer. Neither in metal nor in silicon is positronium formed. Therefore, the positronium formation probability Io-Ps is reduced by the number of positrons penetrating the substrate. In the case of implantation depths below approx. 50 nm, Io-Ps is reduced – the lower the positron implantation depth, the lower is Io-Ps. This reduction of Io-Ps can happen if fewer free electrons are available for the ortho-positronium formation. This is the case for the shorter spur of the positron due to a lower kinetic energy or if free electrons diffuse to the surface. Additionally, this Io-Ps reduction can be interpreted in terms of an out-diffusion of ortho-positronium [19]. In addition to these arguments, which are specific to the PALS method itself, depth-dependent properties in the top 10 nm of the epoxy, such as different densities of electron acceptor groups, could also play their part in the observed decline in Io-Ps. Regarding the average lifetime of ortho-positronium so-Ps for the unaged epoxy film on Au substrate in Fig. 29.2, so-Ps is constant within the experimental scatter for positron implantation depths above approx. 20 nm. This reflects constant free-volume void sizes, independent of the sample depth. A shallow maximum around 150 nm and a decrease in so-Ps above 400 nm are not significant. On the other hand, for very low depths up to 20 nm, a significant increase in so-Ps is observed. This can be interpreted as the effect of a surface region in the order of less than 20 nm wide where the sizes of the free-volume voids increase [19]. However, a diffusion of ortho-positronium to the epoxy surface and into vacuum or an impediment of the positronium formation due to a reduced

Fig. 29.1 PALS analysis of an unaged epoxy thin film on an Au substrate. Ortho-positronium formation probability IoPs is given as a function of average positron implantation depth.

29.3 Results

Fig. 29.2 PALS analysis of an unaged epoxy thin film on an Au substrate. Average ortho-positronium lifetime so-Ps is illustrated with respect to the average positron implantation depth.

spur electron density could have a similar (small) effect on so-Ps in a region very near the surface. 29.3.2 PALS Investigation of Aged Epoxy Films

Fig. 29.3 depicts the effects of aging of the epoxy films on an Au substrate after 100 days of aging. Both aging conditions lead to a significant reduction of Io-Ps. Hygro-thermal aging has a stronger effect than pure thermal aging. The Io-Ps reduction has to be regarded as a hint of an increasing number of electron-acceptor groups. Chemical modifications could be responsible for the Io-Ps reduction. In Section 29.3.3, IR spectroscopy results will be used to give additional information about chemical modifications of the epoxy films due to aging. An Io-Ps reduction is sometimes interpreted in terms of a reduction of the density of free-volume voids in a polymer [10, 20]. However, the influence of the free volume on Io-Ps should be less pronounced than on the ortho-positronium lifetime so-Ps, which depends directly on the average free-volume void size [1]. We will see later that so-Ps varies much less than Io-Ps during aging. Therefore, the Io-Ps reduction should not be interpreted in terms of a free-volume change. After 100 days of exposure, hygro-thermal aging of thin epoxy films on Au leads to an increase in so-Ps of roughly 30–40 ps and hence a bigger free-volume void size than for the unaged sample above average positron implantation depths of approx. 150 nm, while thermal aging reduces so-Ps by roughly 20 ps in the same region. In the near-surface region, i.e., for average implantation depths below approx. 150 nm, so-Ps is reduced for the aged samples in comparison with the values inside the epoxy layer. This has a maximum effect of 30–40 ps for both aging conditions at a depth of approx. 20 nm. The increase in so-Ps for the very low positron

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29 Depth-Resolved Analysis of the Aging Behavior of Epoxy Thin Films by Positron Spectroscopy Fig. 29.3 PALS analysis of the epoxy thin films on an Au substrate. Both ortho-positronium formation probability IoPs and average lifetime so-Ps are depicted as functions of the average positron implantation depth. The data are from an unaged, a thermally aged, and a hygro-thermally aged epoxy film prepared in an identical way. Solid lines are included in the graphs to stress the implantation depth dependence of so-Ps and Io-Ps for every data set.

energy with an implantation depth below 20 nm is similar to that for the unaged adhesive and should have the same reason. Figs. 29.4 and 29.5 show the effects of 100 days’ aging on thin epoxy films on Cu and Al substrates, respectively. The results for Io-Ps and so-Ps for the unaged, thermally aged, and hygro-thermally aged samples are presented. Qualitatively, the epoxy films on Al and Cu show the same behavior as epoxy on Au, but small characteristic differences have to be pointed out. Obviously, Io-Ps is significantly smaller for epoxy films on Cu than for epoxy on Al and Au. Noeske et al. [21] found that 0.1– 0.2 atom% of Cu ions dissolve in the epoxy under the given preparation conditions. This metal ion dissolving effect is negligible for Al and Au. The efficiency of Cu ions in trapping free electrons easily explains the Io-Ps reduction in the case of a Cu substrate (22% for the unaged state instead of 28% on Al and Au). For average positron implantation depths less than approx. 50 nm, Io-Ps decreases in all cases. Whereas so-Ps is in the range 1710 ± 10 ps inside the unaged epoxy films on Al and Au, it is only 1695 ± 10 ps on the Cu substrate. As described below, these films on Cu contain fewer residual oxirane groups than on Al and Au. This could be the reason for the lower size of the free-volume voids in the epoxy network on Cu.

29.3 Results

Fig. 29.4 PALS analysis of the epoxy thin films on a Cu substrate. For details, see Fig. 29.3.

Additionally, both aging mechanisms have less effect on so-Ps in the volume of the epoxy films on a Cu substrate than on Au and Al substrates: for hygrothermal aging there is an increase of 15 ps instead of approximately 30 ps, and for thermal aging the decrease is 20 ps instead of approximately 25 ps. 29.3.3 Further Investigations of Aged Epoxy Films

The curing of the epoxy system is dominated by the crosslinking reaction of DGEBA oxirane and DETA primary and secondary amine groups [22]. Equations (1) and (2) show the addition reactions of DGEBA and DETA. According to infrared spectroscopy results, the oxirane conversion of the unaged thin films is incomplete [16]. Only 65–70% of the oxirane rings are consumed in the adhesives on Au and Al, in comparison with approximately 77% for epoxy on Cu.

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Fig. 29.5 PALS analysis of the epoxy thin films on an Al substrate. For details, see Fig. 29.3.

The higher oxirane conversion for epoxy on Cu is regarded as a catalytic effect of the 0.1–0.2% dissolved Cu ions in the epoxy films [21].

1

2

29.3 Results

It has been shown with IR that the conversion of oxirane groups continues during 100 days of thermal and hygro-thermal aging, resulting in a conversion rate of 80% for the thermally aged epoxy and more than 95% for the hygro-thermally aged epoxy, almost independently of the metal substrate [16]. IR investigations have also given some information about the effects of aging on structural bonds [16]. For the epoxy films on the three metal substrates Cu, Au, and Al, new high-intensity IR bands appear at 1738 cm–1, 1725 cm–1, 1660 cm–1, and 1600 cm–1 due to the aging conditions, proving a remarkable chemical aging in addition to further oxirane conversion. While pure thermal aging mainly leads to formation of an IR band at 1725 cm–1, hygro-thermal aging predominantly enhances IR band formation at 1660 cm–1 and 1600 cm–1. It is concluded that oxidation of remaining reactive amine groups could play a key role in aging [16]. Potential oxidation products such as carbonyl (>C=O) and oxime (>C=N–OH) groups or radical autoxidation mechanisms leading to products such as carbonyl and amide (>N–CR=O) groups could explain some of the IR bands. For hygro-thermal aging, the authors additionally suggest a reaction between carbonyl and residual amine groups leading to imine groups (>C=N–) with an IR band at approx. 1660 cm–1. Fig. 29.6 shows depth-resolved XPS results for the chemical composition of an unaged and a hygro-thermally aged 650 nm epoxy film. The nitrogen atom concentrations are presented with respect to the sputter depth of the adhesive. It is quite constant at 3.9 ± 0.3% through the whole unaged epoxy film. A slight reduction of the nitrogen concentration might be visible at a very low sputter depth of 20 nm. The high nitrogen concentration of approximately 5.0% at a sputter depth of 560 nm could have two reasons: At this sputter depth, part of the epoxy film has been totally removed from the substrate, resulting in a higher measurement inaccuracy due to the lower nitrogen, carbon, and oxygen signals. Additionally, a nitrogen enrichment near the substrate cannot be excluded. In contrast to the unaged film, the nitrogen concentration is only constant at 3.8 ± 0.3 at% for the hygro-thermally aged epoxy at sputter depths above approx. 150 nm (see Fig. 29.7). Near the surface, it is reduced to 3.0 ± 0.3 at%, proving some nitrogen depletion due to hygro-thermal aging. Generally, absolute values

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29 Depth-Resolved Analysis of the Aging Behavior of Epoxy Thin Films by Positron Spectroscopy Fig. 29.6 Depth-resolved XPS results for the analysis of the atomic nitrogen concentrations of an unaged thin epoxy film versus the sputter depth of the adhesive.

Fig. 29.7 Depth-resolved XPS results of the analysis of the atomic nitrogen concentrations of a hygro-thermally aged thin epoxy film versus the sputter depth of the adhesive.

for the nitrogen concentration have to be taken with a pinch of salt due to the possibility of preferential sputtering. In this case, however, the investigation on a 1 mm thick bulk epoxy sample with a known nitrogen content of 4.9 at% reproduced this concentration (measurement: 4.8 ± 0.3 at%) at different sputter depths. Therefore, the concentration data are reliable.

29.4 Discussion and Conclusion

The present investigations show that the combination of PALS, XPS, and IR yields valuable information about the aging behavior of thermosets. For the first time, aging investigations of thin epoxy films have been performed with the PALS beam technique in combination with other experimental techniques. The main quantities for the characterization of polymers with PALS are the average lifetime so-Ps of ortho-positronium and its formation probability

29.4 Discussion and Conclusion

Io-Ps. Special surface effects detected with PALS can be interpreted in terms of the measurement artifacts already mentioned, and need not be regarded as specific surface properties. The unaged epoxy films show depth-independent properties in terms of the free volume, confirming a quite homogeneous average free-volume void size. Additionally, the chemical conversion of the unaged films is far from complete. In this respect, one has to consider the low amount (3.9 ± 0.3 at%) of nitrogen throughout the whole film, since about 4.3 at% of nitrogen corresponds to the stoichiometric ratio of amine hydrogens in the DETA and the oxirane groups in the DGEBA. This cannot explain the low spectroscopic oxirane conversion (approximately 70%) for the unaged epoxies which have already been cured for 48 h at room temperature and for 1 h at 120 8C, however. Hence, additional demixing of DETA and DGEBA is assumed. Both aging conditions, thermal and hygro-thermal aging, reduce Io-Ps significantly. Since the reduction of Io-Ps should be accompanied by an increase in electron acceptor groups, chemical degradation could be the reason for this Io-Ps behavior. The extent of the Io-Ps reduction due to aging is comparable with the Io-Ps reduction in epoxy films on a Cu substrate in comparison with other substrates. There, 0.1–0.2% of dissolved Cu ions in the epoxy matrix are responsible for the Io-Ps reduction. These positive ions are electron acceptors, of course. Hence, it can be concluded that aging produces a similar acceptor group concentration. The Io-Ps reduction is accompanied by complex chemical aging processes indicated by IR investigations. New IR bands suggest the formation of carbonyl, oxime, amide, or imine groups. These groups generally tend to have stronger electron acceptor properties. Thermal aging and (even more so) hygro-thermal aging lead to a further oxirane conversion, starting below 70% for epoxy on Au and Al substrates and at 77% for epoxy on Cu, and reaching 80% after 100 days of thermal aging and over 95% after 100 days of hygro-thermal aging, respectively, for all the metal substrates. This further oxirane conversion should be interpreted in terms of further crosslinking reactions and could have two effects: densification and stiffening of the epoxy network structure. Densification is accompanied by a reduction of free-volume void size (so-Ps PALS measurements). This can be clearly seen for the volume properties of the epoxy films after thermal aging, where so-Ps has decreased for epoxy on Al and Au and a little less for epoxy on Cu, as expected from the stronger additional oxirane conversion due to thermal aging for epoxy on Al and Au than for epoxy on Cu. Since a further oxirane conversion is even stronger during hygro-thermal aging than during thermal aging, one could expect a strong reduction of free volume and therefore of so-Ps in the case of hygro-thermal aging, as well. However, for hygro-thermal aging, an increase in so-Ps and therefore in the free volume is detected. This can be interpreted in terms of swelling of the epoxies, even though they have already dried in vacuum after the 100 days of hygro-thermal aging and most of the water has already escaped according to a desorption experiment. Therefore, it is most probable that the free-volume expansion due

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to residual water during hygro-thermal aging does not relax after removal of the water. IR investigations suggest imine group formation during hygro-thermal aging. This double-bond formation could have a stiffening effect on the epoxy network. Therefore, relaxation processes could be impeded. Both aging regimes show a significant reduction of so-Ps in the top 150 nm of the epoxy film. This corresponds with the nitrogen depletion in this near-surface region detected by XPS after hygro-thermal aging. It is expected that this reflects ammonia or low molecular weight amine desorption resulting in a depth gradient of the nitrogen concentration. A diffusive loss of volatile nitrogen-rich molecules (such as monomeric DETA and small epoxy network fragments) could explain the reduction of so-Ps in this near-surface region if structural relaxation takes place. IR investigations show the formation of new chemical species during both aging processes. However, they cannot definitely prove the formation of volatile low molecular weight and nitrogen-rich molecules, which could escape from the surface region of the epoxy films.

Acknowledgments

The authors thank the German Ministry of Education and Science (BMBF) for financial support of the present investigations.

References 1 Y. C. Jean: Microchem. J. 42 (1990) 72– 2

3

4 5 6

7

8 9

102. F. Faupel, K. Günther-Schade, J. Kanzow, C. Nagel, P. Sperr, G. Kögel: Mater. Sci. Forum 445/446 (2004) 219–223. J. Algers, P. Sperr, W. Egger, G. Kögel, F. H. J. Maurer: Phys. Rev. B 67 (2003) 125 404. O. E. Mogensen: Positron Annihilation in Chemistry, Springer, Berlin (1995). G. Dlubek, H. M. Fretwell, M. A. Alam: Macromolecules 33 (2000) 187. P. Kirkegaard, N. J. Pedersen, M. Eldrup: PATFIT-88, Risø National Laboratory (1989). Y. C. Jean, R. Zhang, H. Cao, J.-P. Yuan, C.-M. Huang, B. Nielsen, P. Asoka-Kumar: Phys. Rev. B 56, 14 (1997) 8459. B. G. Olsen, J. Lin, S. Nazarenko, A. M. Jamieson: Macromolecules 36 (2003) 7618. T. Suzuki, Y. Ito, K. Endo, S. Fujita: Appl. Radiat. Isot. 40 (1989) 397.

10 J. Kanzow, V. Zaporojtchenko, H. Nabi-

11

12

13 14

15 16

17

ka, M. Mizuhata, S. Deki, F. Faupel: Mater. Sci. Forum 445/446 (2004) 313–315. C. Nagel, K. Günther-Schade, D. Fritsch, T. Strunskus, F. Faupel: Macromolecules 35 (2002) 2071. C. Nagel, E. Schmidtke, K. GüntherSchade, D. Hofman, D. Fritsch, T. Strunskus, F. Faupel: Macromolecules 33, 6 (2000) 2242. K. Günther-Schade, D. W. Schubert, F. Faupel: Macromolecules 35 (2002) 9074. K. Günther-Schade, H. L. Castricum, H. J. Ziegler, H. Bakker, F. Faupel: Polym. Eng. Sci. 447 (2004) 1351. C. Wehlack, W. Possart: Macromol. Symp. 205 (2004) 251–261. C. Wehlack, A. Meiser, W. Possart: EURADH, Book of Abstracts for Oral Presentations, DECHEMA (2004) pp. 188– 193. R. Willecke, F. Faupel: Macromolecules 30 (1997) 67.

References 18 A. Thran, T. Strunskus, V. Zaporojtchen-

ko, F. Faupel: Appl. Phys. Lett. 81(2) (2002) 244–246. 19 Y. C. Jean, H. Cao, G. H. Dai, R. Suzuki, T. Ohdaira, Y. Kobayashi, K. Hirata: Appl. Surf. Sci. 116 (1997) 251–255. 20 Y. Y. Wang, H. Nakanishi, Y. C. Jean, T. C. Sandreczki: J. Polym. Sci. B 28 (1990) 1431.

21 M. Noeske, R. Wilken, S. Dieckhoff, O.-D.

Hennemann: (2004) in preparation. 22 J. Kanzow: Struktur, Vernetzung und Ma-

tallisierung duroplastischer Polymere, PhD thesis, University Kiel (2004).

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30 Epoxies on Stainless Steel – Curing and Aging D. Fata, C. Bockenheimer, and W. Possart

Abstract

The chemical aging processes are considered in epoxy adhesive model formulations (diglycidyl ether of bisphenol A (DGEBA) with diethylenetriamine (DETA) for a room temperature (RT)-curing adhesive (EP1) or with dicyandiamide (DDA) for a hot-curing adhesive (EP2)) on a regular stainless steel (Nirosta® 4301; ThyssenKrupp Nirosta) and on a crash-resistant stainless steel (Nirosta® H400; ThyssenKrupp Nirosta). For EP1 layers (3–6 lm), the steel substrates hamper the curing process and hence the glass transition is lower than in the bulk. The effect is more pronounced on Nirosta® 4301. Hot curing provides a fully cured network in EP2 bulk samples. On both substrates, however, the conversion of cyano groups is lower and this leads to less imino ether-like crosslinks and more molecular mobility in the EP2 films. Additionally, a second phase with increased cooperative mobility appears on the substrates just after curing, but not in the bulk samples. Aging is performed in the glassy state at elevated temperatures either in dried air (“thermal aging” – TA) or by immersion in pure water (“hydro-thermal aging” – HTA). Due to the incomplete network in the EP1 layers, curing goes on during both types of aging and changes the microstructure in the layers, but the rate is clearly increased in HTA due to the plasticizing effect of the water that has penetrated. During HTA, a new glass transition with a low Tg is formed, thus indicating the separation of a new amorphous phase. The EP1 network degrades by oxidation of the amines. For TA primary and secondary amines are involved, whereas only the primary amines are oxidized in HTA on both types of stainless steel. All EP2 samples resist chemical degradation in TA. The more mobile phase, which is found in the layers directly after curing, develops in the bulk EP2 only with aging time. Hence, the steel substrates stimulate the observed phase separation. Water induces a similar effect in HTA. Moreover, in all samples, part of the absorbed water is consumed in a chemical reaction that causes the cleavage of the imino ether-like crosslinks. Adhesion – Current Research and Application. Wulff Possart Copyright © 2005 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim ISBN: 3-527-31263-3

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30 Epoxies on Stainless Steel – Curing and Aging

The results show that the crosslinking state and the degradation processes in the epoxy network depend not only on the type of curing agent and the aging regime. The properties of the stainless steel also change these properties within the epoxy interphase at the contact.

30.1 Introduction

Due to the considerable amount of deformation energy that new brands of stainless steel absorb by work hardening, such steel is preferred for making crash absorbers. Adhesive bonding is also indispensable for achieving maximum deformation energy in assemblies of such steel parts. Today, hot-curing epoxy adhesives dominate in automotive production lines, but room temperature (RT)-curing epoxies are needed too, as repair and replacement of hot-curing epoxies by RT-curing adhesives would provide economic advantages. For that reason, the properties of stainless steel joints bonded with epoxy systems are of special interest. According to the literature, the aging behavior of such joints is critical [1, 2]. Mechanical tests reveal that the combined attack of water and temperature causes a strong deterioration in their performance [3]. It is necessary to understand the processes that are going on during aging in order to improve the reliability of structural adhesive bonding. For epoxies in bulk, the following aging mechanisms are reported in the literature. High temperature, which is often denoted as thermal loading, causes not only the less dangerous physical aging of the polymer network [4, 5] but also chemical degradation. The majority of degradation mechanisms attribute breaking of the epoxy backbone to oxygen. For instance, tertiary amines are oxidized to amine oxides, which undergo a Cope elimination and thereby cleave the network [6]. A different mechanism suggests that oxygen radicals attack the carbon atom preferentially in the a-position to the tertiary amine. On recombination of two radicals the epoxy network is scissored [7]. But these mechanisms run only at temperatures higher than 100 8C assisted by UV light, which are uncommon conditions for adhesive joints. In some technical environments as well as for the immersion test in the laboratory, adhesive bonds are exposed to liquid water and elevated temperatures. During that hydro-thermal loading, water diffuses into the epoxy system. It plasticizes the network by shielding attractive physical forces between the polymer chains. In addition, the epoxy network will be chemically altered in some way by the absorbed water. In some cases, water-induced cleavage of network chains is assumed to explain the leaching of low molecular weight products or the loss of mechanical strength. Water-based degradation mechanisms are not proposed, however. Other workers suppose that water opens residual epoxy rings [8–10]. As the results in the literature refer mostly to bulk samples, the question that has to be asked is how aging proceeds in the epoxy/metal interphase, because the joints usually fail in that region.


This paper reports on the aging behavior of an RT-cured and of a hot-cured epoxy network under thermal and hydro-thermal loading in the bulk and in layers on stainless steel. During a period of 100 days, the chemical and structural changes have been observed by IR spectroscopy and differential scanning calorimetry (DSC).


Both of the adhesives used in the experiments contain the diglycidyl ether of bisphenol A (DGEBA) as the bifunctional epoxy resin but they differ in the curing agent. Diethylenetriamine (DETA) as an aliphatic amine hardener with five amine hydrogen atoms is used for the RT-curing adhesive (EP1). The hot-cured single-part epoxy adhesive (EP2) contains dicyandiamide (DDA) as the multifunctional amine-curing agent. No additives are used because we focus on the aging of the epoxy network itself, which represents the main component of any commercial adhesive formulation.

Curing and aging of both epoxy systems is studied on two types of stainless steel. Substrate 1 (SS1) is a common stainless steel with 18% Cr and 8% Ni (Nirosta® 4301). Substrate 2 (SS2) is a stainless steel containing 19% Cr, 7% Mn, and 3% Ni (Nirosta® H400). SS2 performs excellently in crash tests. Both substrates possess rough surfaces due to the treatment by etching and blasting with steel beads (RRMS,SS1 = 2.0 lm, RRMS,SS2 = 2.7 lm).

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30.2.2 Sample Preparation

For EP1, molten DGEBA and DETA are mixed in a mass ratio of 100 : 14 at 55 8C for 5 min followed by fast cooling to RT. Due to the intensive stirring at 55 8C, the mixture is expected to be homogeneous. This is supported by the reproducibility of the curing process at RT (RTC) and by the reproducible network properties. The mixture contains an excess of amine hydrogen as compared with the stoichiometric ratio (DGEBA/DETA = 100 : 12.12). The mixture is cast at 23 8C into silicone molds of appropriate size (for DSC: 3 mm disks, 0.8 mm thick; for IR-ATR (attenuated total reflection) spectroscopy: 10 mm disks, 2 mm thick). For IR-ERAS (external reflection absorption spectroscopy), steel substrates (1 mm thick) are cut to 25 ´ 25 mm2 size. Spin-coating of EP1 (directly after mixing, 10 000 rpm, 10 s) produces thick ca. 5 lm layers on these substrates. For DSC, 7 mm steel disks (0.7 mm thick) are cut and spin-coating provides epoxy layers (ca. 10 lm) on the disks. All samples are first cured at RT for 48 h and subsequently post-cured at Tc = 40 8C for 24 h (PC40). For EP2, DGEBA and DDA are mixed in a mass ratio of 100 : 7 at 55 8C for 30 min followed by degassing and stirring in an ultrasonic bath. This mixing ratio is common for commercial hot-cured epoxy adhesives. The amine hydrogen content is less than the stoichiometric ratio (100 : 13.5). Sedimentation of DDA on curing is avoided by pre-curing at 150 8C for 1.5 h. The resulting prepolymer is solid at RT but dissolves fully in organic solvents such as THF. Hence, no network is formed in the prepolymerization. The prepolymer possesses a low viscosity at 70 8C and no further sedimentation of DDA particles is observed. Bulk specimens of the same size as for EP1 are made by casting the prepolymer melt at 70 8C into Teflon molds. For spin-coating (5000 rpm, 10 s), EP2 is dissolved in THF. This provides layers of ca. 3 lm thickness on the specimens for IR-ERAS and of ca. 6 lm thickness on the disks for DSC. All EP2 samples are heated at about 10 K min–1 from RT to Tc = 180 8C, where the temperature is kept constant for 1 h. Samples are covered by a Teflon foil in order to avoid the evaporation of DGEBA. With all samples, the small epoxy mass produces not much heat of reaction, and the big surface-to-volume ratio avoids any significant overheating on curing. 30.2.3 Aging Experiments

For a physically proper aging simulation, the aging temperature has to be chosen well below the glass transition temperature of the epoxy system. The RTcured EP1 has not a full degree of conversion and it has a relatively low Tg of 55 8C in the bulk. Hence, on thermal aging (TA) at Ta = 40 8C further crosslinking is expected. Therefore, the RT-cured EP1 is post-cured at Ta = 40 8C for 24 h prior to aging. EP2 has a full degree of conversion and a Tg of 177 8C in the bulk after curing. Due to its lower cooperative mobility EP2 is thermally aged at


Ta = 60 8C. Both systems are aged in dried air (dew point < –70 8C; < 400 ppm CO2, 300 mg O2 per liter air). Hydro-thermal aging (HTA) is performed in the glassy state by immersion in distilled water at the same temperatures as for thermal aging. Ta values of 40 8C and 60 8C correspond to 38 mg O2 and 23 mg O2 per liter H2O, respectively. The freshly distilled water possesses 10 lS cm–1 conductivity, which increases with time due to leaching of the epoxy sample and of the glass container. Therefore, the water is exchanged at 20 lS cm–1. 30.2.4 Characterization of Aged Specimens

At selected aging times ta, samples are taken for experimental characterization. After three months of immersion, some samples are re-dried under thermal aging conditions at 40 8C and 60 8C, respectively. With the dry samples, the effect of aging in the epoxy network is studied without artifacts from evaporating water during the experiments. For bulk EP2, DSC is carried out by heating samples (mass ca. 6 mg; PerkinElmer DSC 7) twice from 5 8C to 200 8C at 10 K min–1. Bulk specimens of EP1 and layers of both epoxies on SS1 and SS2 are measured by a MDSC program (modulated DSC, 5–200 8C, average heating rate 10 K min–1, modulation: sinelike, amplitude 0.8 K, period 60 s; TA Instruments Q 100). IR spectroscopy is performed for the bulk on a Bruker IFS 66v/S (EP1) or on a Bio-Rad Excalibur Series FT 3000 spectrometer (EP2) by the attenuated total reflection spectroscopy (ATR) technique (ZnSe hemisphere, p-polarized light, 658 angle of incidence). This technique provides spectral information from a sample surface region of a thickness in the order of the wavelength of the infrared light, i.e., ca. 2.5–12.5 lm. On the steel substrates, the complete epoxy layers are studied by IR-ERAS spectroscopy (p-polarized light, 558 angle of incidence). Both IR techniques make it possible to monitor the consumption of residual epoxy groups and the chemical modifications during aging. In order to make the peak intensities comparable, the spectra are scaled to the intensity of the phenyl band at 1510 cm–1 as the internal standard. In spite of that, the resulting spectra for the bulk and for the layers may not be compared quantitatively because of the different spectroscopic techniques used. IR microscopy is performed in the ATR mode both on the bulk and on the epoxy–steel samples made for the DSC.

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30.3 Results and Discussion 30.3.1 The RT Curing Epoxy System (EP1) 30.3.1.1 Curing of EP1 The curing reactions start immediately upon mixing of DGEBA and DETA. A primary or a secondary aliphatic amine group couples with an oxirane group by ring opening and a hydroxide group is formed at the link between the monomers. Bulk samples and spin-coated layers are prepared in the initial stage of these polyaddition reactions. The spectra do not show any reaction of OH groups with oxirane rings for the catalyst-free EP1 because of the low Tc and the much more reactive amine groups. Corresponding to the epoxy conversion, the glass transition temperature Tg increases (see Fig. 30.1). In the bulk, EP1 vitrifies slowly between 211 and 271 min after mixing (MDSC) when the glass transition temperature Tg exceeds RT. Hence, the reaction rate slows down because further curing is diffusion-controlled. As soon as the system is in the glassy state an endothermic peak is superimposed on the high-temperature side of the glass transition and increases with time (Fig. 30.1). This peak indicates physical aging in the glassy network, which is no longer in thermodynamic equilibrium. The nonequilibrium state causes a slow relaxation of enthalpy.

Fig. 30.1 MDSC characterization of EP1 at different curing times at RT (RT = vertical line).


With increasing difference (Tg–Ta) and shorter aging time ta, the amount and the rate of enthalpy relaxation diminish. The enthalpy relaxation is accompanied by structural changes that are described as “short-range order” (SRO). (This well-accepted term may be misleading because this structure relaxation is usually not associated with a measurable structural order. Experiments indicated that the structural changes could be described as a local rearrangement of polymer segments that results in a reduced mean distance between the segments.) Further details on enthalpy relaxation and physical aging of polymers are given in Refs. [4], [11], and [12]. The EP1 layers vitrify later (at 230–325 min), indicating a hampered curing on SS1 and on SS2 as well. An aging peak on Tg similar to that for the bulk shows that physical aging is present (Fig. 30.2). In the following, we use the temperature Tmax at the peak maximum as a convenient measure for the glass transition. After RT curing (RTC), only &67% of the epoxy groups are consumed in the bulk (l-IR-ATR, Fig. 30.3) and Tbulk max = 55.5 ± 0.1 8C (DSC, Fig. 30.4) after 48 h. RT curing leads to lower values in EP1 on SS1 (Tlayer max = 47.1 ± 0.8 8C) and on SS2 (Tlayer = 49.8 ± 0.5 8C). Obviously, the cooperative macromolecular mobility at RT max is higher in the glassy 10 lm RTC layers than in the bulk and it depends on the type of stainless steel. The low Tg values and the correspondingly reduced degrees of conversion in the layers (Fig. 30.3) confirm that RT curing is hampered on the steels. Above Tmax, an exothermic peak appears in the first DSC run (Fig. 30.2). It represents the heat of thermally stimulated post-curing in the viscoelastic poly-

Fig. 30.2 MDSC characterization of the RTC–EP1 in the bulk on SS1 and SS2.

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Fig. 30.3 Tmax of the aging peak (DSC) at the glass transition of EP1 in the bulk and in 10 lm layers on SS1 and SS2 in the RTC and in the PC40 states.

Fig. 30.4 IR intensity of the normalized epoxy band (915 cm–1, l-ATR) for EP1 in the bulk and in 5 lm layers on SS1 and SS2 for the RTC and for the PC40 states.


mer. The IR-ATR spectra reveal no residual epoxy groups after the first DSC run. Hence, this thermal treatment provides full epoxy conversion in the bulk. The final glass transition appears at Tbulk = 129.0 ± 0.2 8C for the EP1 bulk in a g second DSC run. On both substrates, however, the post-curing peak is considerably weaker than for the bulk EP1 (Fig. 30.2). The second DSC run provides no further curing and a final Tlayer of only 93.4 ± 6.0 8C on SS1 and 91.9 ± 8.0 8C on g SS2. Correspondingly, the IR microscopy reveals a considerable number of residual epoxy groups in the epoxy layers after the first DSC run. These data show that EP1 does not leave the glassy state when RTC samples are post-cured at 40 8C for 1 h (PC40), either in the bulk or in the layers. Therefore, PC40 proceeds under diffusion control and is only little accelerated. As for RTC, the gross reaction rate slows down due to the decreasing mobility and only some additional oxirane conversion is obtained, but not full consumption (Fig. 30.3). Hence, the glass transition rises to only Tbulk max = 66.6 ± 0.3 8C in the bulk PC40 (Fig. 30.4 and the curve for 0 days in Fig. 30.6, below). The smaller aging peak at the glass transition results from the higher crosslink density achieved by PC40. For the layers, PC40 leads to a similar rise in Tg (Tlaxer max = 57.2 ± 1.5 8C layer on SS1; Tlayer = 60.6 ± 2.6 8C on SS2) but the T values remain lower than Tbulk max max max (Fig. 30.4). As expected, the degrees of oxirane conversion also tend to increase (Fig. 30.3) in the layers, but remain below the bulk values.

30.3.1.2 Thermal Aging of EP1 after Post-Curing at 40 8C During TA at 40 8C, the glass transitions shift to higher temperatures in all EP1–PC40 samples (Figs. 30.5–30.8). After 100 days, we get Tbulk max = 80.1 ± 0.1 8C while Tlayer rises to 75.7 ± 2.0 8C (on SS1) and 76.7 ± 0.8 8C (on SS2) (Fig. 30.5). max The accompanying consumption of epoxy groups is verified by the decay of the IR band at 915 cm–1. Obviously, TA stimulates some further curing in EP1– PC40 at low rates in the bulk and in the layers on both stainless steels. For the bulk, the exothermic post-curing peak appears in the first DSC run (Fig. 30.6). Surprisingly, no post-curing peak is observed any longer for the layers after two days’ TA (Figs. 30.7 and 30.8) although residual epoxy and amine groups are present. Something must have happened in the EP1–PC40 layers on the two steels that prevents the usual increase in the chemical reaction rate caused by raising the temperature. It is proposed that the hampered degree of conversion results from a substrate-induced separation of an amine-rich phase and an epoxy-rich phase. On aging, the spatial separation of the reactants additionally retards their diffusion-controlled curing in the glassy network. Obviously, even the fast heating into the viscoelastic state cannot stimulate sufficient interdiffusion of the reactants close to SS1 and SS2. For the bulk EP1–PC40, the “aging peak” first decreases for short TA and then increases again (Fig. 30.6). The decrease is explained in the following way. Tg and hence (Tg–Ta) increase at a high rate due to intensified epoxy consumption in the early stages of TA with Ta = constant. Now, the amount of relaxed free volume decreases in a network with growing difference (Tg–Ta). Conse-

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Fig. 30.5 DSC: Tmax(aging peak) of EP1–PC40 in the bulk and on SS1 and SS2 during thermal aging at 40 8C up to 100 days.

Fig. 30.6 DSC curves for EP1–PC40 in the bulk during thermal aging at T = 40 8C up to 101 days.


Fig. 30.7 DSC for EP1–PC40 on SS1 during thermal aging at T = 40 8C up to 100 days.

Fig. 30.8 DSC for EP1–PC40 on SS2 during thermal aging at T = 40 8C up to 100 days.

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quently, old SRO regions dissolve at a higher rate than new SRO regions are generated, and this results in a decreasing aging peak. For longer aging times, (Tg–Ta) becomes almost constant because Tg does not rise any more. Hence, more short-range ordering develops than is dissolved, leading to a growing aging peak. The layers behave differently again. The aging peak grows continuously (Figs. 30.7 and 30.8) even though the glass transition temperature increases as for the bulk. Therefore, layers and bulk possess different microstructures after TA. When bulk EP1–PC40 is aged at 40 8C, two new peaks show up in the IRATR spectra at 1660 cm–1 and at 1730 cm–1 (Fig. 30.9). They are attributed to a product 1 and a product 2, respectively. IR-microscopy on the fracture surface of a thermally aged bulk specimen provides no such peaks, however. It is concluded that only some of the surface region undergoes chemical modifications under the influence of oxygen and temperature on the time scale of our experiments. The modifications are proposed to proceed as oxidation of the remaining primary and secondary amine groups on exposure to air. A primary amine is converted into an oxime group, which is identified as product 1 [Eq. (1)]. Its concentration evolves on a low level with ta (Fig. 30.10). Oxidation of a secondary amine splits the polymer chain and produces an oxime and a carbonyl group (products 2 and 3, Eq. (2)). Hence, this second mechanism degrades the network. Such a thermal degradation is expected to decrease Tg [13] and both oxidation reactions impair the ongoing cure as they also consume amines. Neither a drop in Tg nor a loss of mechanical strength can substantiated so far by our

Fig. 30.9 IR-ATR: spectral region (1800–1560 cm–1) during thermal aging of EP1–PC40 in the bulk at 40 8C up to 112 days.


Fig. 30.10 IR-ATR: intensity of the normalized aging band at 1660 cm–1 for EP1– PC40 in the bulk during thermal and hydro-thermal aging at 40 8C up to 85 days.

data, however, since product 3 appears in very low concentration and the oxidation is confined to a thin surface region (IR-ATR information depth *10 lm), which is not relevant for the DSC curves of bulk samples. R2

n HC = R1

R2 NH2 O2

R1

R2

n 2 HC = R1

!

N H

R3 = C H 3O2 n R4

n CN =

OH H2 O

1

1

R2 n ! 2 C O 2HO = R1 2

R3 = N C 2H2 O n R4 3

2

In the EP1–PC40 layers on SS1 and SS2, the same products 1–3 evolve during TA. Now, their concentrations depend on the kind of steel substrate (Fig. 30.11). SS1 stimulates the oxidation of primary amines (to product 1) more than SS2. (Note that the data in Fig. 30.11 may not be compared quantitatively with those in Fig. 30.10 because the IR spectroscopic techniques are different.)

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Fig. 30.11 IR-ERAS: intensity of the normalized aging band at 1660 cm–1 for EP1–PC40 on SS1 and SS2 during thermal and hydro-thermal aging at 40 8C up to 85 days.

30.3.1.3 Hydro-thermal Aging of EP1 After one day of HTA at 40 8C, the bulk EP1–PC40 absorbs 1.7 wt.% of water. Over 16 days, the water content increases rapidly to 4.5 wt.% and reaches 5.07 ± 0.03 wt.% after 100 days. Fig. 30.12 summarizes the DSC results during HTA. The Tbulk max (first glass transition in Fig. 30.12) drops from 66.6 ± 0.2 8C to 62.8 ± 0.2 8C within the first day. This increased cooperative mobility in the network reveals the plasticizing effect of the absorbed water. During further HTA, the intensity of the exothermal post-curing peak decreases steadily (DSC curves, Fig. 30.12) and this is accompanied by a relatively fast consumption of oxirane rings (IR-ATR spectroscopy, peak at 915 cm–1 depicted in Fig. 30.13). After 25 days of HTA almost full consumption of epoxy groups is reached. Simultaneously, Tbulk max shifts upward to 73.9 ± 0.2 8C within 100 days (Fig. 30.9), indicating a loss of cooperative mobility in the hydro-thermally aged EP1–PC40 due to the ongoing cure. (The time scales of the IR and the DSC results are not comparable because of their different information depths and the different sample geometry.) In addition, samples which are re-dried after 100 days of HTA have their Tbulk at 129.2 ± 0.3 8C. g Hence, it is concluded that during HTA at 40 8C an almost completely crosslinked network develops in the bulk. Thermal aging of EP1–PC40 (see Section 30.3.1.2) over the same period of time results only in a slight increase in the degree of epoxy conversion and a weak rise of Tbulk (80.1 ± 0.1 8C). Obviously, the g water in the network increases the rate of ongoing curing significantly. This is explained by two effects. First, the plasticizing effect of water promotes the dif-


Fig. 30.12 DSC measurement of EP1–PC40 in the bulk during hydro-thermal aging at T = 40 8C up to 101 days.

Fig. 30.13 IR-ATR: intensity of the normalized epoxy band (915 cm–1) for EP1– PC40 in the bulk during thermal and hydro-thermal aging.

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fusion-controlled reaction by giving an increased molecular mobility. Second, water catalyzes the curing reaction of amine with epoxy groups according to the literature [14, 15]. After 24 days of HTA at 40 8C, a second glass transition appears at Tbulk g2 & 88 8C, well above the first glass transition (Fig. 30.12) at Tbulk g1 & 71 8C. Obviously, a phase with lower cooperative mobility forms in the plasticized bulk matrix of EP1–PC40 during HTA. Hydro-thermal aging of EP1–PC40 bulk samples produces only one new IR band, at 1660 cm–1. It is assumed that this band belongs to the same product 1 (with oxime functionality) discussed for thermal aging. For long aging times, the concentration of product 1 is similar for TA and HTA (see Fig. 30.10). Since no band develops at 1730 cm–1 or anywhere else in the spectra, it is concluded that the presence of water suppresses the oxidation of secondary amines in the bulk samples. When hydro-thermally aged samples of EP1–PC40 layers on steel are taken from the immersion bath their big surface/volume ratio leads to a quick re-drying at RT. Therefore, the first DSC run characterizes layers with less water content than in the bath. During the first run (not shown), a broad endothermic peak masks all the other heat flows, especially the Tg of the epoxy layer. This endothermic peak corresponds to the heat of desorption of the residual water. MDSC makes it possible to separate the heat of desorption (nonreversing signal) from the heat flow at a glass transition (reversing signal). The corresponding Tlayer g, rev. for the EP1–PC40 layers increases with aging time, as depicted in Fig. 30.14. Although Tg,? is not reached within 100 days, the rising Tlayer g, rev. val-

Fig. 30.14 MDSC: Tgrev of EP1–PC40 on SS1 and SS2 during hydro-thermal aging up to 100 days.


ues indicate a water-stimulated ongoing cure similar to that of the bulk. The corresponding epoxy consumption is verified by IR-ERAS (not displayed). However, a considerable number of epoxy groups remain unreacted in the network after 100 days’ HTA. Concerning chemical modification, only the band at 1660 cm–1 arises in the epoxy layers, indicating the formation of oxime groups (product 1) in analogy to the bulk. In the layers, the oxime concentration is greater for HTA than for TA (Fig. 30.11). Obviously, the presence of water enhances the oxidation of primary amines in the layers. Furthermore, the oxime concentration depends on the type of steel: For HTA, it is lower on SS1 than on SS2. This is just the opposite of the relationship found for TA (Fig. 30.11). 30.3.2 The Hot-Curing Epoxy System (EP2) 30.3.2.1 Curing of EP2 The solid DDA powder is insoluble in the epoxy resin at low temperatures. Additionally, the tautomeric structure of the DDA molecule, shown schematically in the Introduction, stabilizes the amine groups. For these reasons, mixtures of DGEBA and DDA do not react at room temperature. Curing starts when hardener molecules diffuse into the liquid epoxy and this occurs well below the melting point of DDA at 213 8C. In the diffusion zone three types of chemical crosslinks are formed by the three curing reactions: Type 1 crosslinks result from the addition of primary and secondary amines to oxirane groups, as shown for EP1 (Scheme 30.1 a). The type 2 crosslinks are due to the etherification of an epoxy with a hydroxyl group, which proceeds well at the elevated curing temperature (Scheme 30.1 b). The reaction of a cyano group with an OH group forms a type 3 crosslink, being specific for DDA-cured systems Scheme 30.1 c). Adjacent to the imino ether, the type 3 crosslinks include a new secondary amine that is capable of forming an additional type 1 crosslink with residual epoxy groups. A second OH group adds slowly to the imino ether, generating an ester and another secondary amine (Scheme 30.1 d). This reaction does not change the number of crosslinks and is not completed at the end of curing because the IR spectra show the band of the imino ether (1680 cm–1) as well as that of the ester (1730 cm–1). The cured EP2 bulk has its Tbulk at 177.4 8C. The IR spectroscopy shows g 100% conversion of epoxy groups, type 1 and type 2 crosslinks, an almost full conversion of cyano groups and type 3 crosslinks. On both types of stainless steel, Tlayer reaches 162 8C after curing (Fig. 30.15). This is significantly lower g than in the bulk, even though the epoxy conversion is 100% in the layers as well. The reason is that the cyano conversion is slightly lower in the layers than for the bulk (l-IR-ATR, cyano doublet around 2210 cm–1; Fig. 30.16). As a result, the cured network of EP2 on SS1 and SS2 possesses a slightly increased cooperative mobility due to the less abundant type 3 crosslinks. Both for the bulk and the layers, curing is not hampered by vitrification during network formation because the Tg values are lower than the Tc (180 8C) at

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(a)

(b)

(c)

(d)

Scheme 30.1 The curing reaction in EP2: formation of (a) type 1 crosslinks; (b) type 2 crosslinks; (c) type 3 crosslinks; (d) additional type 1 crosslink.

any time. The epoxy does not vitrify before it is cooled, and therefore no physical aging takes place during the cure. On a microscopic scale, the curing reactions are influenced by the diffusion of DDA molecules from the solid DDA particles into the molten DGEBA. Due to the diffusion, the concentration of hardener decreases with distance from the DDA grains. This concentration gradient prevails even in the cured network because more and more DDA molecules are fixed by chemical bonds to the growing network. Consequently, there is an inhomogeneous distribution of crosslink density according to the distribution of DDA particles and the DDA concentration profile in the mixture. Such inhomogeneities in latent curing adhesives are discussed in the literature [16, 17] and verified there by TEM, small-angle X-ray scattering and small-angle neutron scattering.


Fig. 30.15 Glass transition temperature of EP2 in the bulk and in the layers on SS1 and SS2 after curing.

Fig. 30.16 l-IR-ATR: spectral region (2300–2100 cm–1) of EP2 in the bulk and in layers on SS1 and SS2 after curing.

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30.3.2.2 Thermal Aging of EP2 For both bulk samples and EP2-layers on SS1 and SS2, no significant chemical changes are detected in the IR spectra during TA at 60 8C. The constant intensities for the cyano band (2210 cm–1) and the epoxy band (915 cm–1) reveal no further crosslinking. In the bulk, the principal glass transition temperature Tbulk g1 remains at the initial value of 177 8C (DSC; Fig. 30.17) during TA. After 1 day of TA, a second glass transition appears at Tbulk g2 (ta = 1 d) = 86 8C, indicating the formation of a new amorphous phase in the inhomogeneous glassy network. It becomes more pronounced on further aging and Tbulk g2 has shifted to 114 8C after 108 days. That phase separation is reversible because the second glass transition vanishes as the aged sample is heated to the viscoelastic state and characterized again (second DSC run). For the EP2 layers on SS1 and SS2, the glass transition of the matrix stays constant at Tlayer g1 = 162 8C on TA (Figs. 30.18 and 30.19). Surprisingly, a second glass transition is found at Tlayer g2 = 105 8C immediately after curing, very visible on SS2 at least. The second glass transition is still there in a second DSC heating run. Therefore, a second thermally stable amorphous phase has formed on curing in the layers. On TA, Tlayer does not shift significantly but the height of g2 the Tg step tends to increase slightly. After heating above Tlayer g1 , the second phase is still present at Tlayer but with a reduced step height. It is assumed that g2 EP2 undergoes a phase separation in the layers which is different from the bulk. On SS1 and SS2 the new phase is not fully dissolved in the viscoelastic

Fig. 30.17 DSC measurement of EP2 in the bulk during thermal aging at T = 60 8C up to 108 days.


Fig. 30.18 DSC measurement of EP2 on SS1 during thermal aging at T = 60 8C up to 100 days.

Fig. 30.19 DSC measurement of EP2 on SS2 during thermal aging at T = 60 8C up to 100 days.

499

500


state and Tlayer is significantly higher than Tbulk g2 g2 (ta = 1 d). Consequently, the microstructure of EP2 close to SS1 and SS2 is different from the microstructure of the bulk.

30.3.2.3 Hydro-thermal Aging of EP2 The water uptake of EP2 in the bulk reaches 2.6 wt.% after one day of HTA at 60 8C and increases rapidly to 4.2 wt.% over 15 days. Afterward, the rate of water uptake slows down and at 99 days 5.17 ± 0.05 wt.% is reached. Corresponding with the main part of water uptake Tlayer g1 , decreases from 177.4 8C initially to about 105 8C within 35 days (Fig. 30.20) due to the plasticizing effect of water. Then the principal glass transition stays at this temperature. The plasticization shifts Tbulk by –72 K, which is very high in comparison with g1 shifts of –20 K to –30 K reported for other hot-cured epoxies [18, 19]. According to gravimetric analysis, about 1 wt.% of water remains in the network after the first DSC run and this amount is still present after the second run. Other time–temperature regimes (e.g., 48 h at 105 8C) do not remove this residual water content either. Therefore, it is concluded that 1 wt.% of water is bonded irreversibly to the bulk network. Tbulk g1, re-dried reaches 134.1 ± 0.3 8C after 99 days’ HTA, which is well below Tbulk g1 = 177 8C of the non-aged network. Therefore, the total plasticization effect (maximum shift DTbulk g1 = –72 K within 99 days) splits into an irreversible and a reversible part. The irreversible part

Fig. 30.20 DSC measurement of EP2 in the bulk during hydrothermal aging at T = 60 8C up to 99 days.

30.3 Results and Discussion bulk corresponds to the difference (Tbulk g1 –Tg1, re-dried), which amounts to 32 K at one day and 43 K at 99 days of HTA. See below for the underlying mechanism. Again, a second glass transition develops below the principal glass transition during HTA. The height of this cp step and an endothermic aging peak increase, and its characteristic temperature Tbulk max2 rises from 90 8C to 99 8C with ongoing aging (Fig. 30.17). Due to the plasticizing effect of water, the final Tbulk max2 is about 15 K lower than for TA. The second glass transition is more distinct in the case of HTA because Ta = 60 8C is closer to Tbulk g1 in the plasticized epoxy system. As a result, there is more macromolecular mobility in the network and this improves phase separation. The aging peak that appears at the upper end of the second glass transition after eight days indicates enhanced local packing due to physical aging in the second phase. After a second DSC run, the sample shows only the principal glass transition at Tbulk g1 for the modified matrix network. The separation of a second bulk phase is also reversible for HTA. The layers of EP2 on SS1 and SS2 are at least partially re-dried shortly after removal from the immersion bath due to their high surface/volume ratio. As for EP1, the first DSC run of the layers does not characterize the same state as for the bulk, with almost full water content. But the first run removes the remaining free water as for the bulk, and therefore the second runs for the layers and for the bulk are comparable. Tlayer (second run) decreases on SS1 from 161.7 ± 2.3 8C to 137.3 ± 2.4 8C and on g1 SS2 from 160.5 ± 2.2 8C to 135.2 ± 0.9 8C (Fig. 30.21). Hence, there is an irreversible plasticization in the EP2 layers on both types of stainless steel, as for the bulk. Especially after a long exposure to water, the Tlayer (second run) values are similar g1 for the layers and the bulk even though the initial values are significantly different. It is assumed that the irreversible plasticization is caused by the breaking of bonds under the influence of water. This hydrolytic aging process should reduce the crosslink density of the epoxy network. The hydrolysis does not deteriorate the entire network, however, because the irreversible plasticization comes to an end within the period of HTA we investigated. The IR spectroscopic results support the assumed partial hydrolysis of the EP2 network. Both for the bulk and the layers, a new IR band grows at 1730 cm–1 during hydro-thermal aging while the imino ether band at 1680 cm–1 decreases (e.g., Fig. 30.22 for EP2 on SS2). These changes are similar to the addition of OH groups to the imino ether observed on curing. But the chemical modification on HTA is induced by water since it is not observed during TA. Hence, we propose the degradation mechanism depicted in Scheme 30.2 for the bulk and for the layers. The imino ethers are hydrolyzed by water and an ester group, assigned to 1730 cm–1, and a primary amine are formed. The reaction is well known in organic chemistry as the splitting of a Schiff’s base and runs at moderate temperatures [20]. For the epoxy network, breaking the crosslinks increases the polymer mobility and this explains the irreversible decrease in Tg1 during hydro-thermal aging. The fact that Tg1, re-dried reaches a plateau at long aging times indicates that only the type 3 crosslinks are hydrolyzed. The crosslinks of types 1 and 2 prove to be stable under the applied hydro-thermal load.

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502


Fig. 30.21 Tg(second run) of EP2 in the bulk and in the layers on stainless steel on hydro-thermal aging at 60 8C up to 100 days.

Fig. 30.22 l-IR-ATR: spectral region (1850–1550 cm–1) of EP2 on SS2 during hydro-thermal aging at 60 8C up to 100 days.

30.4 Conclusion

Scheme 30.2 The hydrolysis mechanism proposed for EP2 in the bulk and in the layers on hydro-thermal aging.

This prevents the total disintegration of the network. The plateau reached for the bulk and for the layers is similar because curing leads both in the bulk and in the layers to a similar density of type 1 and type 2 crosslinks because of the full oxirane conversion. This confirms the assumption that the lower Tg1 of EP2 in the layers on SS1 and SS2 after curing results from a lower concentration of type 3 crosslinks as compared with the bulk. During the contact with water all type 3 crosslinks are hydrolyzed in the bulk and in the layers at a similar rate. The hydrolysis mechanism is also confirmed by the fact that 1 wt.% of water is irreversibly absorbed into the network. This corresponds to a water content of 0.055 mol of water per 100 g of the epoxy system, which matches well the concentration of imino ether groups in the fully cured network (0.078 mol per 100 g). The proposed hydrolysis mechanism differs from a similar mechanism mentioned in the literature for an epoxy system containing DDA [21]. There, water which has to be present during curing is considered to prevent the formation of type 3 crosslinks. Accordingly, the reaction between cyano and OH groups is replaced by the reaction of cyano with water, thus already forming a carbonyl group and a primary amine during curing.

30.4 Conclusion

A RT-curing (EP1) and a hot-curing epoxy adhesive (EP2) prepared as bulk materials and as layers on two types of stainless steel have been investigated during curing and aging in dried air and in distilled water. After RT curing of EP1 the macromolecular mobility in the layers is significantly higher than in the bulk. In addition, the mobility of EP1 on SS1 is higher than on EP1 on SS2. Parallel to the difference in mobility, more reactive groups remain in the layers on SS1 than on SS2. In the bulk, the lowest number of reactive groups remain. This ranking of the substrates and the bulk prevails on further curing. Consequently, the stainless steel substrates hamper the curing of EP1 in a way specific to the type of stainless steel, probably due to a

503

504


substrate-induced phase separation. Both in the bulk and in the network of the layers, the existence of short-range ordered regions is verified by enthalpy relaxation during RT curing and post-curing at 40 8C as well. On thermal aging at 40 8C, the curing reaction proceeds in the bulk and in the layers at a similar rate. The ongoing reaction reduces the mobility in the glassy network more and more. This leads to an asymptotic approach of the degree of conversion. The initial content of short-range ordered regions diminishes for the bulk but increases for the layers. Consequently, thermal aging leads to different microstructures in the bulk and in the layers. On hydro-thermal aging, water diffuses into the epoxy system and plasticizes the network. The resulting mobility increases and the catalytic effect of water promotes the ongoing curing. In the bulk, the full degree of conversion and the lowest mobility is achieved for EP1, whereas a higher mobility and a considerable number of reactive groups remain in the layers. A second glass transition appears during hydro-thermal aging of EP1, at least in the bulk, and verifies the existence of an inhomogeneous network with at least two amorphous phases. In the contact zone with air, the network degrades by the oxidation of primary and secondary amines. In the presence of water, only primary amines are oxidized. Consequently, the amines are consumed by the ongoing cure and by oxidation. These mechanisms complete the picture of the degradation mechanisms discussed in the literature for low aging temperatures. Hydro-thermal aging at the presence of both types of stainless steel increases the oxidative modification in EP1. Curing of EP2 provides a fully cured network in the bulk and a relatively low cooperative mobility (Tg = 177 8C). The molecular mobility of EP2 is similar on SS1 and SS2 but significantly lower than in the bulk due to a lower degree of cyano conversion. Correspondingly, the layers contain a lower density of imino ether-like crosslinks while the densities of amine- and ether-like crosslinks are like those in the bulk. Only in EP2 on SS1 and SS2 is a second phase with increased mobility present just after curing. On thermal aging, no chemical changes are detected either in the bulk or in the layers. In the bulk, thermal aging gives rise to a more mobile phase, whereas the second phase in the layers stays at the initial temperature found for it after curing. Therefore, the microstructure of the EP2 networks in the layers differs from the microstructure in the bulk. Water induces a similar phase separation at least in the bulk. It breaks the imino ether-like crosslinks specific for the DDA-cured epoxy systems in the bulk as well as in the layers on the two types of stainless steel. The proposed mechanism describes the aging behavior of DDA-cured epoxy systems well, since it explains all the observed effects of chemical modification, irreversible plasticization, and irreversible water uptake. Additionally, it is understood why there is no complete disintegration of the network: the hydrolysis cleaves only the imino ether-like crosslinks but not the amine-like or the ether-like crosslinks that are also formed during curing. Hence, after hydro-thermal aging the macromolecular mobility in EP2 is similar in the layers and in the bulk because the content of amine-like and ether-like crosslinks is similar.

References

In summary, the chemical degradation mechanisms of the epoxy systems are determined by the curing agent, the type of aging, and, occasionally, the combination of epoxy system and steel substrate.

Acknowledgment

The financial support of the Studiengesellschaft Stahlanwendung e.V., Düsseldorf, is gratefully acknowledged (Project P553).

References 1 O. D. Henneman, R. Lüschen, Adhäsion 2

3 4

5

6 7 8

9

10

1992, 36, 37–41. R. D. Adams, J. Comyn, W. C. Wake, Effect of the environment on structural adhesives, in Structural Adhesive Joints in Engineering, Chapman & Hall, London, 1997, Chap. 8. K. H. Daniel, 41th International SAMPE Symposium 1996, pp. 1773–1781. L. C. E. Struik, Physical Aging in Amorphous Polymers and other Materials, Elsevier Scientific Publishing, Amsterdam, 1978. M. J. Richardson, Curing of thermosets, in Calorimetry and Thermal Analysis of Polymers (Ed.: V. B. F. Mathot), Munich, 1994, Chap. 7. B. L. Burton, J. Appl. Polym. Sci. 1993, 47, 1821–1837. V. Bellenger, J. Verdu, J. Appl. Polym. Sci. 1985, 30, 363–374. G. Z. Xiao, M. E. R. Shanahan, J. Polym. Sci.: Part B: Polym. Phys. 1997, 35, 2659– 2670. A. Raveh, D. Marouani, R. Ydgar, J. E. Klemberg-Sapieha, A. Bettelheim, J. Adhes. 1991, 109–124. G. M. McMaster, D. S. Soane, IEEE Trans. Components, Hybrids, Manuf. Technol. 1989, 12(3), 373–386.

11 D. J. Plazek, J. Polym. Sci.: Part B: Polym.

Phys. 1990, 28, 431–448. 12 C. Bockenheimer, D. Fata, W. Possart, J.

Appl. Polym. Sci. 2004, 91, 361–368. 13 L. Barral, J. Cano, A. J. López, J. Lopez,

14 15 16

17 18 19

20

21

P. Nógueira, C. Ramírez, Thermochim. Acta 1995, 269/270, 253–259. L. Shechter, J. Wynstra, Ind. Eng. Chem. Res. 1956, 48, 94. I. T. Smith, Polymer 1961, 2, 95. R. Schmidt, Inhomogenitäten bei amorphen Polymeren, in Polymere Werkstoffe; Vol. I Chemie und Physik (Ed.: H. Batzer), Georg Thieme Verlag, Stuttgart, 1985, Chap. 1.6. U. T. Kreibich, R. Schmid, J. Polym. Sci.: Symposium 1975, 53, 177–185. D. M. Brewis, J. Comyn, J. A. Shalash, J. L. Tegg, Polymer 1980, 21, 357–360. R. A. Pethrick, E. H. Hollins, I. McEwan, A. J. Mackinnon, D. Hayword, L. A. Cannon, S. D. Jenkins, P. T. Mc Grail, Macromolecules 1996, 29, 5208–5214. F. Möller, Amine durch Spaltung von Schiffschen Basen, in Houben-Weyl: Methoden der Organischen Chemie; Vol. XI/1: Stickstoffverbindungen II, 2003, Chap. VIII i). M. D. Gilbert, N. S. Schneider, Macromolecules 1991, 24, 360–369.

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31 Scanning Kelvin Probe Studies of Ion Transport and De-adhesion Processes at Polymer/Metal Interfaces K. Wapner and G. Grundmeier

Abstract

Depending on the environment of an adhesive joint, moisture or electrochemical/corrosive reactions at the interface which lead to a chemical degradation of bonds could lead to a failure of the joint. In a corrosive environment, electrochemical reactions dominate, in particular on reactive metals at room temperature, whereas the electrode potential characterizes the reactivity of the interface. The scanning Kelvin probe, which measures the Volta potential difference between a specimen and the calibrated sensing probe, is introduced as the only electrochemical technique which allows nondestructive, real-time measurements of electrode potentials at adhesive/metal oxide interfaces in situ, even if they are covered with an adhesive layer. A complementary view is presented of the adhesion and de-adhesion mechanisms, especially in humid and corrosive environments, which are predominant and most important for the application of metal/adhesive composites in engineering applications. The transport of hydrated ions at metal/adhesive interfaces is considered as an important premise for corrosive reactions. A height-regulated scanning Kelvin probe (HR-SKP) which is capable of measuring topography and interfacial electrode potential at the same time has been introduced for the study of rough or shaped surfaces. This new HR-SKP has been applied on different samples of interest, focusing on measurements in situ while in a state of active corrosion. It has been shown that this system is suitable for measurements of active cathodic delamination or filiform corrosion of organically coated iron or aluminum specimens. A combination of a pressurized blister test and the HR-SKP can be utilized for the study of adhesives under a combined corrosive and mechanical load.


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31 Scanning Kelvin Probe Studies of Ion Transport and De-adhesion Processes

31.1 Introduction

Adhesive joins are becoming an increasingly interesting technology for the automotive industry. The use of high strength steels, combinations of materials, or pre-primed metal sheets promotes the usage of adhesives. High strength joints with excellent initial properties are already available. Still, there is a lack of understanding of the fundamentals of the adhesion forces and the mechanisms that lead to a weakening of interfacial bonds with time of exposure in corrosive environments. Depending on the environment of the joint, moisture or electrochemical/corrosive reactions at the interface which lead to a chemical degradation of bonds could lead to a failure of the joint. In corrosive environments, electrochemical reactions dominate, in particular at room temperature and on reactive metals. Electrochemical reactions require: · the establishment of an electrified interface · the conservation of charge, which in particular requires a balance between cathodic (with electrons crossing the interface from the substrate) and anodic reactions (with electrons crossing the interface to the substrate or positively charged ions crossing the interface from the substrate). Electrified interfaces are predominantly built up by mobile charged species which in the case of metal/polymer interfaces may be identified with ions embedded in the polymeric matrix. Therefore, electrochemically driven reactions prevail in environments which allow the presence of such species. Furthermore, the conservation of charge may require the presence of ion-transfer reactions and this condition is almost always satisfied for reactive metals such as iron, copper, zinc, and aluminum. As the nature of the electrified interface dominates the kinetics of corrosive reactions, it is most desirable to measure, e.g., the drop in electrical potential across the interface, even where the interface is buried beneath a polymer layer and is therefore not accessible for conventional electrochemical techniques. The scanning Kelvin probe (SKP), which measures in principle the Volta potential difference (or contact potential difference) between the sample and a sensing probe (which may consist of a sharp wire composed of a conducting, stable phase such as graphite or gold) by the vibrating condenser method, is the only technique which allows the measurement of such data and therefore all modern models which deal with electrochemical de-adhesion reactions are based on such techniques [1–8]. Recently, it has been applied mainly for the measurement of electrode potentials at polymer/metal interfaces, especially polymercoated metals such as iron, zinc, and aluminum alloys [9–15]. The principal features of a scanning Kelvin probe for corrosion studies are shown in Fig. 31.1.

31.2 Theory and Experimental Set-Up of a Scanning Kelvin Probe

Fig. 31.1 Schematic representation of the measurement of electrochemical potentials at buried interfaces (here cathodic delamination of an adhesive layer from an iron substrate) with the scanning Kelvin probe. Volta potential difference DW, Kelvin current i(t) induced by the vibration of the needle (reference), compensation voltage U0.

31.2 Theory and Experimental Set-Up of a Scanning Kelvin Probe

For the simple case of two electronic conductors in vacuum, the Kelvin probe measures the work function difference between a sample and the probe using the vibrating capacitor method [11, 16–18]. The Kelvin probe, which may consist of a sharp tip of a wire composed of a conducting, stable phase like graphite or gold, is connected electrically to the sample under investigation. This leads to identical Fermi levels in both the sample and the probe and, in the case where the work functions of the two materials differ, to a charging of one phase with respect to the other (giving a Volta potential difference or contact potential difference, DW). Vibration of the probe establishes an alternating current i(t) between the sample and the probe forming a charged capacitor. For the measurement of the Volta potential difference, a compensation voltage U0 which is equal to –DW could be added to the electric circuit; thus the alternating current is compensated to zero (nulling technique) (Eq. (1), with dQ(t)/dt the alteration of the charge, U0 the compensation voltage applied between the two electrodes, dC(t)/dt the variation of the capacitance, and t the time). it

dQt dCt DW U0 dt dt

1

For a parallel plate capacitor with the sinusoidally modulated plate separation d(t) given by Eq. (2) (d0 = average distance between the parallel plates, Dd = vibration amplitude and x = angular frequency of vibration), the time dependent capacitance C(t) is given by Eq. (3) with C(t) defined by Eq. (3).

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31 Scanning Kelvin Probe Studies of Ion Transport and De-adhesion Processes

dt d0 Dd cos xt Ct

2

eA eA dt d0 Dd cos xt

3

Using a small, flattened needle as the scanning probe, Baumgärtner et al. [19] suggested a ratio between the needle diameter and its average distance of at least 5 : 1 to ensure that an assumed homogeneous field distribution of a parallel plate capacitor is valid. Combining Eqs. (1) and (3) to obtain the time-dependent Kelvin current i(t) gives Eqs. (4) and (5). it DW U0 eA DW U0

d 1 dt d0 Dd cos xt

exADd sin xt d0 Dd cos xt2

4 5

Assuming that the average distance between the parallel plates is much greater than the vibration amplitude (Dd 1. Since this was not found among a variety of adhesive joints, the influences of combined cyclic loading and high temperature cannot be separated. From the viewpoint of polymer physics, this is to be expected since the segmental motion afforded by higher temperatures increases the sensitivity of the polymer to the mean stress in cyclic loading – creep rupture may occur prior to fatigue. As a consequence, S–N curves have to be determined at each relevant temperature. Static load which causes creep of adhesive joints requires time-dependent knock-down factors. With a requirement of maximum shear strain of 0.1 after 10 years, some flexible adhesives already require extremely low factors (0.001– 0.07) at ambient temperature. Therefore, stresses due to static long-term loading in flexible adhesive joints have to be kept low or should even be avoided. Attempts to express the knock-down factor of combined temperature and static long-term loading by the product of the factors of individual influences yielded a result similar to combined temperature and cyclic loading. Again, this is a question of time scales of segmental motion at different temperatures. With the experimental data and available methods of finite element analysis, a part of a side wall with a window was designed, manufactured, and tested (Fig. 33.9). This part represents a modular structural component of a coach. A rigid adhesive system (Araldite® 2015) was used to bond the load-carrying steel parts, which were pretreated using CLP. The GFRP panels and the window were bonded to the structure using a flexible adhesive (Terostat MS937). Global–local modeling and verification techniques were used to identify and to assess the critical areas in the structure. Table 33.6 shows the results of a typical fatigue loading test at ambient temperature, which was carried out on our structure. The test is passed if the struc-

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33 Adhesive Joints for Modular Components in Railway Applications Fig. 33.9 Part of a side wall of a rail vehicle.

ture does not break after having seen all the cycles at given stress amplitudes. The first two columns contain the load amplitudes Fa and mean loads Fm applied in the test. In the next two columns, stress amplitudes ra and mean stresses rm calculated in the most critical area of the stiff adhesive bond line are given. For this calculation, a finite element model of the whole structure and a local model of the bond line was used. Stresses were calculated using the material model given in Section 33.4.1. Column 5 shows stress amplitudes which follow by transforming columns 3 and 4 to a stress ratio of R = –1 using Eq. (8). The corresponding predicted cycle numbers for 95% survival probability calculated from Eq. (6) are given in column 6. Finally, the number of cycles which were applied to the structure in the fatigue test appear in column 7. A comparison of columns 6 and 7 shows that the cycle number applied is always less than the theoretical cycle number for 95% survival probability. Therefore,

Table 33.6 Comparison of a fatigue life prediction for a rigid, load-carrying adhesive joint (Araldite 2015) and fatigue life test results for part of a side wall.

Fa [kN]

Fm [kN]

ra [MPa]

rm [MPa]

ra,–1 [MPa]

Npredicted

Ntested

30.0 25.0 20.0 13.5 10.0 2.0 1.2

0.0 0.0 0.0 16.5 15.0 3.0 1.8

4.50 3.75 3.00 2.03 1.50 0.30 0.18

0.00 0.00 0.00 2.48 2.25 0.45 0.27

4.50 3.75 3.00 2.33 1.70 0.31 0.18

6 ´ 104 2 ´ 105 1 ´ 106 6 ´ 106 6 ´ 107 >1010 >1010

1 1 13 1 ´ 106 1 ´ 106 1 ´ 106 2 ´ 106

Acknowledgment

following the prediction, the structure should not fail. In fact, the structure could not be broken in the fatigue loading test. Although this result shows that the assumptions are conservative, it is not possible to draw conclusions on the validity of the underlying models since fatigue failure was not achieved. More experimental work is needed.

33.6 Conclusion

In this paper, we have presented test results of two methods of adhesion improvement on a specific aluminum alloy and a stainless steel. The methods are based on flame pyrolysis of silane precursors and a laser pretreatment in connection with a primer. While the first method seems to be most favorable for steel surfaces, the latter gives good results on both aluminum and steel. Both methods are well suited for application in the railroad industry since they are effective under atmospheric conditions and are therefore promising for integration in an automated production line. Because of the complex geometry of parts and components containing adhesive joints, finite element models are generally very large. Hence, pragmatic solutions are needed for the description of the adhesive properties. As a consequence, besides short-term properties only specific aspects of long-term properties such as creep and fatigue can be handled. Approaches for the determination of material parameters have been discussed and data for some commercial adhesives were presented. It was shown that cyclic loading may cause a change in the failure mechanism leading to a reduced fatigue life. In cases where adhesive failure is observed under static loading conditions, fatigue life appears to be very short. Therefore, it should be emphasized that testing of adhesive joints under cyclic loading conditions is mandatory for every qualification program. Adhesion problems may then be identified and conclusions on the surface pretreatment can be drawn. Testing under different stress ratios indicated that transformation rules for low mean stresses seem to exist. The dependence of the segmental motion on time scales in polymers leads to a breakdown of conventional design rules. Special attention is required in cases where certain conditions appear simultaneously. Especially at temperatures near and above the glass transition, strength reduction due to multiplication of partial knock-down factors might be underestimated. Testing under relevant conditions is needed instead. Therefore, the most important issue in adhesive joint design is a critical consideration of the loading conditions in service life with respect to the corresponding temperatures and time scales.

Acknowledgment

Financial support by the BMBF is gratefully acknowledged (03N3084A/B).

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33 Adhesive Joints for Modular Components in Railway Applications

References 1 UIC Kodex 566 VE, 3rd edn., Internatio2

3

4 5

6

7

naler Eisenbahnverband, Germany, 1990. A. S. Argon, in: Materials Science and Technology, Vol. 6, R. W. Cahn, P. Haasen, E. J. Kramer, Eds., VCH, Weinheim, 1993. B. Crist, in: Materials Science and Technology, Vol. 12, R. W. Cahn, P. Haasen, E. J. Kramer, Eds., VCH, Weinheim, 1993. Huntsman, WO 9803600 (1998) and WO 9623037 (1996). R. D. Adams, J. Comyn, W. C. Wakel, Structural Adhesive Joints in Engineering, 2nd edn., Chapman & Hall, London, 1997. A. J. Kinloch, Adhesion and Adhesives: Science and Technology, Chapman & Hall, London, 1990. A. Wulf, M. Brede, O.-D. Hennemann, Numerical simulation of thick adhesive

8 9

10 11 12 13

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joints: Application of the theory of hyperelastic continua, Proc. Euradh 98, 1998. R. W. Ogden, Proc. R. Soc. London A 1972, 326, 565–584. I. M. Ward, D.W. Hadley, An Introduction to the Mechanical Properties of Solid Polymers, Wiley, New York, 1993, Chap. 10. B. P. Haigh, J. Inst. Metals 1917, 18, 55– 86. B. P. Haigh, Engineering (London) 1917, 104, 315–319. D. Radaj, Ermüdungsfestigkeit, Springer, Berlin, 1995. A. Schöpfel, H. Idelberger, D. Schütz, D. Flade, Materialprüfung 1997, 39(6), 246– 251. M. H. Beheshty, B. Harris, Comp. Sci. Technol. 1998, 58, 9–18.

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34 Behavior of Dismantlable Adhesives Including Thermally Expansive Microcapsules Y. Nishiyama and C. Sato

Abstract

A novel technique for dismantling joints bonded with an adhesive is presented and its behavior is described. An epoxy adhesive including thermally expansive microcapsules, which consist of plastic shells filled with liquid hydrocarbon and expand at high temperature, has been developed. The volume of the adhesive, including the microcapsules, can also increase due to the expansion of the microcapsules even though the adhesive has been fully cured. Therefore, adhesive joints can be dismantled easily by increasing the temperature. Since the total expansion of the adhesive depends on the content of the microcapsules, its expansion was measured for different contents of microcapsules. Several kinds of experiments were carried out on joints bonded with the adhesives containing different weight fractions of microcapsules, by immersing the joints in boiling water or by keeping them in hot air. Cracks were generated and then propagated mainly at the interfaces between the adhesive and the adherend due to thermally caused internal stress. The expansion of the adhesive and the thermally expansive microcapsules was also measured at several temperatures and pressures using a pressure–volume–temperature (PVT) apparatus. The maximum volume expansion of the adhesive under atmospheric pressure was about 400% at 90 8C. In this case, the microcapsules in the adhesive expanded to about 800%. The microcapsules could transfer a pressure of only 1–2 MPa to the matrix resin according to the result of the PVT test. That pressure was not sufficient to deform the matrix resin at room temperature. Therefore, both the expansion of the microcapsules and the softening of the matrix resin must be achieved for the expansion of the adhesive.


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34 Behavior of Dismantlable Adhesives Including Thermally Expansive Microcapsules

34.1 Introduction

Recycling of industrially produced material is important from the environmental point of view. For the recycling of materials bonded with adhesives, dismantlement of the products is necessary to separate them, although that is quite difficult. Therefore, a novel technique by which adherends can be dismantled easily has to be established. A conventional way to solve the problem is the use of thermoplastic adhesives because they soften at high temperature and joints bonded with them can be separated by applying small forces. Recently, a new kind of adhesive including thermally expansive microcapsules has been proposed as an alternative method [1]. A joint bonded with such an adhesive can be dismantled by heating, as shown in Fig. 34.1. Originally, the microcapsules were used as an expanding and pressurizing agent for the molding of composite materials in sandwich panels [2]. Another special application of the microcapsules was as fillers of removable coatings used in car painting lines [3]. However, applying the material to adhesives is quite a new challenge. The first research on such adhesives was conducted by Sakurai et al., in order to bond plywood boards [4]. They used silicone-modified epoxy resin as the matrix of the adhesive and added small amounts of the thermally expansive microcapsules to it. Cohesive fracture of the adhesive due to the expansion of the microcapsules at high temperature was reported. Ishikawa modified and improved the technique for bonding wallpaper on plywood or plasterboards in construction fields [5–7]. He used quite soft adhesives, e.g., vinyl copolymers, including the thermally expansive microcapsules, because high strength was not required for these applications. The layer of adhesive delaminated interfacially and the metal

Fig. 34.1 Image and schematic illustration of the dismantlable adhesive filled with thermally expansive microcapsules.


sheets could be peeled off very easily by hand. We extended this novel technique to structural adhesives by changing the adhesive from such a soft material to the harder and stronger epoxy resin. The behavior of the adhesive system has been investigated in terms of the composition [8–11]. In this paper, we present the mechanical aspects of the novel adhesive. At first, experiments using the adhesive were carried out to investigate the dismantlability and the joint strength as a function of the content of the microcapsules. Next, the volume change of fully cured bulk adhesive was measured with respect to its temperature and the weight fraction of the microcapsules. The dismantlability of the joints was evaluated by observing them after heating. Lap shear tests of the joints were also carried out to measure the bond strength. The expansion behavior of the microcapsules was measured by a PVT (pressure–volume–temperature) apparatus [12, 13] since it is also very important for the design or for modification of the microcapsules. The expansion behavior of the adhesive is discussed on the basis of the experimental results.


As the filler of the dismantlable adhesive, thermally expansive microcapsules (Matsumoto Microsphere F-30; Matsumoto Yushi-Seiyaku Co. Ltd.) were used. They consisted of a poly(vinyl chloride) shell filled with liquid hydrocarbon (liquefied isobutane) as shown in Fig. 34.2. The diameter of the microcapsule was 10–30 lm and the average was 20 lm. The shell thickness was 3–4 lm at room

Fig. 34.2 Structure of the thermally expansive particles.

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Fig. 34.3 Chemical structures of Epikote 828 and Epomate B002 (n = degree of polymerization).

temperature. As the temperature rose, the internal pressure of the shell increased due to the vaporization of the liquid hydrocarbon and the stiffness of the shell decreased. As a result, the volume of the microcapsules expanded as shown in the images in Fig. 34.2. The microcapsules began to expand drastically at 60–70 8C and their volume finally increased 50–100-fold. Bisphenol A type epoxy resin (Epikote 828; Shell) cured with modified amine (Epomate B002; Shell) was used as the matrix of the dismantlable adhesive (resin/matrix = 2 : 1 w/w). Fig. 34.3 shows their chemical structures. The bulk adhesive was cured at room temperature for 24 h before the experiments. Cured bulk resin mixed with the microcapsules was used for the specimens to measure the volume change. 34.2.2 Volume Expansion of the Cured Bulk Adhesive

The thermal expansion ratio of the adhesive containing the thermally expansive microcapsules was measured with a thermodilatometer as shown in Fig. 34.4. The syringe was made of glass, so the specimen could be observed by eye. The bulk cylinders (diameter: 7 mm, length: 15 mm) of the cured adhesive were

Fig. 34.4 The thermodilatometer.


used as specimens in the experiment. The specimen was placed in a syringe filled with silicone oil and it was heated by a heater through an outer oil jacket. The volume change of the specimens containing different weight fractions of microcapsules up to 50% was measured isothermally at different temperatures. The same experiment was carried out using only silicone oil and its volume expansion was substracted from the experimental result. The volume change of the specimens was calculated from the translation of the piston measured with a displacement pickup. 34.2.3 Dismantlability of Joints Bonded with the Dismantlable Adhesive

The dismantlability of joints bonded with the adhesive was evaluated experimentally. Fig. 34.5 shows the configuration and the dimensions of such joints. Specimens of three sizes (20 mm ´ 20 mm, 50 mm ´ 50 mm and 100 mm ´ 100 mm, all with 0.3 mm bond line thickness) were prepared. The adherends of the specimens were made of aluminum alloy (5052H34, Si: £ 0.25%, Fe: £ 0.40%, Cu: £ 0.10%, Mn: £ 0.10%, Mg: 2.2–2.8%, Cr: 0.15–0.35, Zn: £ 0.10, Al: rest). The aluminum surfaces were finished with two types of surface pretreatment. For half of the adherends, the surfaces were polished using fine grinding paper (#2000, SiO2). For the other half of the adherends, the surfaces were milled using rough grit-blasting (#80, SiO2). Before bonding, the adherends were degreased with acetone. They were bonded with the adhesives containing 0–50 wt.% of microcapsules. The bonded specimens were cured at room temperature for 24 h. To dismantle the specimens, they were immersed in boiling water or held in hot air at 100 8C for up to 2 h and the situation of the specimen was observed visually. 34.2.4 Bond Strength of the Dismantlable Adhesive

Lap shear tests were carried out using single lap joint specimens of the aluminum alloy (5052H34). Sample dimensions are shown in Fig. 34.6. Surface pretreatment and curing conditions of the joint specimens were similar to the other specimens used for the dismantlement test. The lap shear strength of the

Fig. 34.5 Configurations and dimensions of joint specimens bonded with dismantlable adhesive.

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34 Behavior of Dismantlable Adhesives Including Thermally Expansive Microcapsules Fig. 34.6 Configuration and dimensions of a single lap joint specimen.

specimens was examined using a mechanical tensile tester at room temperature. Crosshead speed was 0.1 mm s–1. 34.2.5 PVT (Pressure–Volume–Temperature) Tests

The thermal expansion of the microcapsules was measured at several temperatures under several hydrostatic pressures using a PVT (pressure–volume–temperature) apparatus which was made to measure the volume change of materials [14]. In the PVT apparatus (Fig. 34.7), the pressure and the temperature could be controlled independently and the volume change of the specimen could be measured as a function of pressure and temperature.

Fig. 34.7 Schematic illustration of the PVT apparatus.


The specimen was enclosed in a bellows vessel with silicone oil. This vessel was placed in a pressure vessel filled with another silicone oil and it was heated by an electric heater through the pressure vessel and outer oil jacket. The outer silicone oil was pressurized by a piston and an actuator. Thus at a given temperature, we could apply a hydrostatic pressure to the specimen through the bellows vessel. The volume change of the specimen was calculated from the displacement of the bellows measured using an LVDT (linear variable differential transformer). Inside the transformer, a pin was connected to an iron core and the pin was pushed by the bellows. At a translation of the bellows, the core, which was pressurized in a stainless steel pipe, also moved and the position of the core was detected by an LVDT coil from the outside of the pipe. The experiment with the microcapsules was carried out again using a syringe vessel. In this case, only the specimen was enclosed in the syringe, without any silicone oil, so that the change of total volume including the microcapsules and leaked gas could be measured. Because the PVT tests started after pressurization, the gas expanded under a set pressure. The volume fraction of the microcapsules in the syringe was 75%.

34.3 Results and Discussion 34.3.1 Volume Expansion of the Cured Bulk Adhesive

Fig. 34.8 shows the volume expansion of the specimens. With the temperature of the silicone oil near the specimen as the parameter, the volume changes are plotted as a function of the weight fraction of the microcapsules. The expansion of the unfilled adhesive was less than that of the epoxy filled with microcapsules. The volume increase starts at 60 8C and almost levels around 110 8C. Obviously, the final volume expansion depends on the weight fraction of the microcapsules. In the case of 50 wt.% of microcapsules, the final expansion was

Fig. 34.8 Volume expansion of adhesives containing thermally expansive particles at several temperatures as a function of their weight fraction.

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34 Behavior of Dismantlable Adhesives Including Thermally Expansive Microcapsules Fig. 34.9 Volume change of cured bulk adhesive including 50 wt.% of thermally expansive microcapsules (left: unheated; right: heated to 120 8C).

over 400% in volume. Fig. 34.9 shows a bulk resin containing 50 wt.% of microcapsules. The smaller sample consists of room temperature (RT)-cured resin prior to exposure to heat. The larger sample is the same resin after additional heating to 120 8C for about 30 s. Although the matrix resin was fully cured and very hard, the bulk specimen expanded very greatly. 34.3.2 Dismantlability of Joints Bonded with the Dismantlable Adhesive

The adhesive layer of several specimens expanded and peeled off from both of the interfaces to the adherends when they were heated in either boiling water or hot air. However, the other specimens having only a small content of the thermally expansive microcapsules did not show such peeling. We defined three types of dismantlement of the specimens: total dismantlement, protrusion of adhesive out of the joint, and partial dismantlement. They are categorized by visual inspection. Total dismantlement means spontaneous separation of the adhesive layer from the adherends and there was no necessity to apply any force to separate them. Since no spots of adhesive residue could be observed by eye on the surface of the broken interfaces in this case, interfacial fracture had occurred. In the specimens showing protrusion of the adhesive, the adherends were connected by the adhesive layer and the separation of the adherends by hand was difficult, although the adhesive layer around the rim of the specimen debonded interfacially. Partial dismantlement of the joints denotes the condition in which the adhesive layer was not totally separated from the adherend at the interfaces but was connected by several bridges. The joints were quite easily separated by hand. Therefore, they can be used in terms of adherend recycling, as shown in Fig. 34.10. Tables 34.1 and 34.2 show the results of the experiments for the specimens with the grinding paper finish and for the grit-blasted specimens, respectively. According to these results, the dismantlement of the specimens became easier with increasing weight fraction of the thermally expansive microcapsules, and immersion in boiling water was more effective than keeping them in hot air. The specimens


Fig. 34.10 Specimens after heating: (a) total dismantlement; (b) adhesive protrusion; (c) partial dismantlement.

Table 34.1 Results of dismantlement tests on specimens with a grinding paper finish. a)

Specimen size [mm2]

Weight fraction [%] Hot water

20 ´ 20 50 ´ 50 100 ´ 100 a)

Hot air

0

10

20

30

40

50

0

10

20

30

40

50

UC UC UC

UC UC UC

SO SO UC

TD TD TD

TD TD TD

TD TD TD

UC UC UC

UC UC UC

UC UC UC

UC UC UC

UC PD PD

TD PD TD

TD: totally dismantlement; SO: sticking out of adhesive; PD: partially dismantlement; UC: unchanged.

finished with grinding paper delaminated more easily than the samples with a grit-blast finish. In other words, the specimens having smoother interfaces showed better results in terms of dismantlability from a macroscopic viewpoint. It has been mentioned previously that cohesive fracture occurred when a soft resin was used as the matrix of such dismantlable adhesives [5]. However, with

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34 Behavior of Dismantlable Adhesives Including Thermally Expansive Microcapsules Table 34.2 Results of dismantlement tests on specimens with a grit-blasted finish.a)

Specimen size [mm2]

Weight fraction [%] Hot water

20 ´ 20 50 ´ 50 100 ´ 100 a)

Hot air

0

10

20

30

40

50

0

10

20

30

40

50

UC UC UC

UC UC UC

UC UC UC

SO SO UC

SO SO SO

TD TD SO

UC UC UC

UC UC UC

UC UC UC

UC UC UC

UC UC UC

SO UC SO

TD: totally dismantlement; SO: sticking out of adhesive; PD: partially dismantlement; UC: unchanged.

a hard and high-strength resin as the matrix of the adhesive, it had dismantlability and interfacial fracture was observed. Therefore, the mechanism of dismantlement of the hard and high-strength adhesive seems to be quite different from that of the soft adhesive. In the case of the soft resins, the strength of the resins is so low that cohesive fracture in the resin phase could easily occur. On the other hand, using the hard and strong resin, cohesive fracture cannot occur because its strength is much greater than the interfacial strength between the adhesive and the adherends. Thus, interfacial fracture takes place due to the internal stress caused by the expansion of the adhesives. 34.3.3 Bond Strength of the Dismantlable Adhesive

The fracture surfaces of the specimens after the test were different, as shown in Fig. 34.11, and they depended on the surface treatment. Fig. 34.11 a and b show

Fig. 34.11 Fracture surfaces of single lap joints: (a) with a grinding paper finish; (b) with a grit-blasted finish.

34.3 Results and Discussion Fig. 34.12 Variation of failure strength of single lap joints with respect to the weight fraction of thermally expansive particles.

Grit-blasted finish Grinding paper finish

typical fracture surfaces for the specimens with grinding paper finish and with grit-blasted finish, respectively. The grinding paper finished specimens delaminated interfacially. However, the grit-blasted specimens always showed cohesive fracture. Fig. 34.12 shows the lap shear strength of the joints. The lap shear strength increased with respect to the weight fraction of the thermally expansive microcapsules when it was lower than 30 wt.%. The strength a showed maximum (12.3 MPa for grit-blasted finish and 7 MPa for grinding paper finish) around 30 wt.% of the microcapsule content and decreased after the peak. The specimens finished with grinding paper were much stronger than those with gritblasting, so surface roughness must be very influential on the strength. 34.3.4 PVT Relationship of Microcapsules and Dismantlable Adhesive

Fig. 34.13 shows the thermal volume expansion of the microcapsules with the hydrostatic pressure as the parameter, obtained by using the bellows vessel and a heating rate of 3 K min–1. For confirmation of the heating rate, the same results were obtained when the PVT test was carried out using 2 K min–1 as the heating rate. The microcapsules start to expand dramatically around 60 8C both

Fig. 34.13 Variation of thermal volume expansion of microcapsules, with pressure as a parameter, measured using the bellows vessel.

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34 Behavior of Dismantlable Adhesives Including Thermally Expansive Microcapsules Fig. 34.14 Variation of thermal volume expansion of microcapsules, with pressure as a parameter, measured using the syringe.

at atmospheric pressure (0.1 MPa) and at 1 MPa. However, the expansion under 2 MPa and 3 MPa started around 80–90 8C. The volume of the capsules swelled by a factor of more than 60 at 100 8C under atmospheric pressure but not more than 10 times at 1 MPa. However, the volume under high pressure decreased after passing a peak. In this result, it seems that the hydrocarbon in the microcapsules was leaked out and dissolved in the silicone oil under high pressure. The expansion behavior of the microcapsules was also measured by using the syringe (Fig. 34.14). In this case, the volumes not only of the microcapsules but also of the leaked hydrocarbon could be measured together. Hence, the volume increased continuously. Fig. 34.15 shows the thermal expansion of the RT-cured adhesives under several pressures measured using the bellows vessel. The adhesives expanded greatly, even though they had been cured. The maximum volume expansion under atmospheric pressure was about four-fold at 90 8C. Above 100 8C, it seems that volume increased again. However, the volume expansion was caused by leaked gas; it was confirmed that the gas leaked from the specimen by thermodilatometry with the glass syringe (Fig. 34.4). Since the microcapsules in the adhesive were expanded about eight-fold, only 1–2 MPa of hydrostatic pressure was applied from the microcapsules to the matrix resin, as shown by the intersection in Fig. 34.14.

Fig. 34.15 Variation of thermal volume expansion of dismantlable adhesive containing 50 wt.% microcapsules, with pressure as a parameter, measured using the bellows vessel.

34.5 Conclusion

34.3.5 Discussion

Dismantlability of the joints depends on both the macroscopic roughness of the aluminum adherends and the weight fraction of the microcapsules. The lap shear strength also depends on these conditions. Great roughness of the adherends was not advantageous for increasing the dismantlability, but it was an advantage for obtaining higher bond strength. Therefore, it can be said that the roughness and strength are in a relationship of compromise. The weight fraction of the microcapsules produces a similar situation. The lap shear strength increased with respect to the weight fraction of the thermally expansive microcapsules when it was less than 30 wt.%. The strength showed a maximum (12.3 MPa in the grit-blast finishing condition and 7 MPa for grinding paper finishing) around 30 wt.% of the microcapsules and then began to decrease. The specimens finished with grinding paper were much stronger than those grit-blasted, so surface roughness must be very influential on the strength. When we applied the adhesive with 50 wt.% of microcapsules to the adherends finished by grit-blasting (80), the highest strength (9.68 MPa) of the joints was obtained though the adhesive was totally dismantled in boiling water. Even though the strength was lower than that of recent high-performance structural adhesives, we might say that the dismantlable adhesive has enough strength for semi-structural applications. The expansion behavior of the microcapsules observed using the syringe method seems to be quite similar to the situation inside the dismantlable adhesive, because the hydrocarbon gas cannot leak from the matrix resin so it accumulates. On the other hand, since the hydrocarbon from the microcapsules near the adhesive surface can leak out easily, the expansion behavior of the microcapsules near the adhesive surface can be simulated using the bellows vessel method.

34.5 Conclusion

We have conducted a feasibility study on dismantlable adhesive joints for structural use. Even though a rigid epoxy resin was used, the joints bonded with 30 wt.% of thermally expansive microcapsules in the adhesive could be totally dismantled. In that case, interfacial fracture occurred. The lap shear strength of the adhesive was quite high and it exceeded 10 MPa under advantageous conditions. The strength depends strongly not only on the weight fraction of the microcapsules but also on the macroscopic roughness of the adherend surfaces. The dismantlability of the adhesive joints also depends on both of these conditions. Even under high pressure, the microcapsules can swell. However, the changes in pressure and volume were not sufficient to rupture the matrix resin. Therefore, both the expansion of the microcapsules and the softening of the matrix resin are indispensable for the expansion of the adhesive.

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References 1 Y. Nishiyama, N. Uto, C. Sato and H. Sa-

10 C. Sato, Y. Nishiyama, M. Kobayashi and

kurai, Int. J. Adhes. Adhes., 2003, 23, 377–382. US 5 242 637; EP 0 407 996. US 4 844 833; JP 2 131 217 and 2 136 376. H. Sakurai et al., Bulletin of Shizuoka Industrial Technology Research Center, 1998, 43, 11–16 (in Japanese). H. Ishikawa, Proc. JSME Colloquium, 2001, 01-86, 5–8 (in Japanese). H. Ishikawa et al., J. Adhes. Soc. Japan, 2004, 40, 146–151 (in Japanese). H. Ishikawa et al., J. Adhes. Soc. Japan, 2004, 40, 184–190 (in Japanese). M. A. Hanafi, Y. Nishiyama, N. Uto and C. Sato, Extended Abstracts of EURADH 2002, 2002, 133–136. Y. Nishiyama and C. Sato, Proc. 8th Japan International SAMPE Symposium, 2003, 661–664.

T. Koto, Proc. 17th International Symposium SWISSBONDING 2003, 2003, 467– 468. Y. Nishiyama, C. Sato, N. Uto and H. Ishikawa, J. Adhes. Soc. Japan, 2004, 40, 298–304 (in Japanese). C. Sato, Y. Nishiyama and N. Uto, Proc. 27th Annual Meeting of The Adhesion Society, Inc., 2004, 256–258. Y. Nishiyama and C. Sato, Extended Abstracts of EURADH 2004, 2004, 1, 319– 324. P. Zoller, P. Bolli, V. Pahud, and H. Ackermann, Rev. Sci. Instrum., 1976, 47, 948.

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569

Subject Index a adherence energy 62 adhesion – autohesion 389, 393 – bioadhesion 175 – cell 159 – chemical 4, 5, 11, 24, 265, 282 – covalent bonds 5 – diffusion polymer-polymer 382, 387 – durability 217 – force measurement, atomic force microscope (AFM) 33 – interdiffusion layers 196 – interfacial width 382 – polymer/polymer 196 – protein 157 – thermodynamic work 33 adhesion force, viscoelastic contribution 33 adhesion promoters, hyperbranched polymers 217 adhesive 217 – biocompatible polymers 193 – controlled polymerization 190 – controlled superstructure 190 – development 190 – dismantlable 555 – flexible 540 – high-performance 189 – hot-melt pressure-sensitive adhesive 229 – isocyanate-free polyurethane chemistry 193 – molecular design 190 – multicomponent multiphase 190 – nanofillers 194 – pressure-sensitive adhesive 229 – programmed biosynthesis 192 – rigid 540 adhesive joint – design 550 – design limits 525

– dismantlable 555 – durability 540, 550 – elastic 525, 529 – mass transport application 525 – material behavior 319, 540 – mechanical properties 403 – nondestructive testing (NDT) 403 – nonlinear ultrasonic testing 403 – modeling 531, 540 – railway applications 539 – stiff 528 – stress distribution 528, 529 – structural analysis 525, 540 – testing 525 adsorption 63 – from liquid phase 3 – isotherm 4 – poly(methyl methacrylate) 47 – preferential 72 AFM s. atomic force microscopy Ag 89 aging – cataplasma storage 541 – chemical 445 – crosslinking during 445 – degradation 479 – durability 445 – effects 72 – epoxy networks 446 – hydro-thermal 479 – hygro-thermal 466 – hygro-thermo-oxidative 447 – phase separation 479 – physical 445 – thermal 447, 466, 479, 487 – thermo-oxidative 447 aluminum (Al) 1, 72, 89, 133, 137, 138, 146, 217, 265, 407, 408, 409, 410, 412, 445, 466, 518, 541, 559


570

Subject Index amine, diffusion of diethylenetriamine (DETA) 135 aramid 221 atomic force microscopy (AFM) 33, 47, 59, 342 – contact mode 33 – force distance curve 37 ATR s. attenuated total reflection spectroscopy attenuated total reflection spectroscopy (ATR) 74 Au 72, 138, 445, 466

b barnacle adhesive – adhesive properties 143 – aluminum 143 – polydimethylsiloxane (PDMS) 143 – structure formation 143 – supramolecular structure 143 benzophenone 252 bioadhesion 176 bioadhesives 150 biocolloids 152 biocompatibility 47 biodegradability 47 biofilm, bioadhesion 176 blister test 507 block copolymer – atomic force microscopy (AFM) 342 – polystyrene/polyisoprene 341 – small angle scattering (SAXS) 342 – styrenic 337 BM s. Brillouin microscopy bond 2 – covalent 2 – van der Waals 2 bonding, interfacial 2 Brillouin microscopy (BM) 125, 126, 132 – sound attenuation 133 – sound velocity 133 Brillouin spectroscopy (BS) 126

CLSM s. confocal laser scanning microscopy compatibilizer 199 composite materials – carbon/epoxy 307 – glass/epoxy 307 – hyperbranched polymers (HBP) 217 – surface treatment 305 confocal laser scanning microscopy (CLSM) 163 contact angle measurement 272, 289 contact mechanics – contact radius 379 – energy release rate 379 – semicrystalline polymer-glassy polymer 365 continuum mechanics – adhesive joints 319 – Cosserat model 319 – interphase 319 copolymer blends – elastic modulus 233 – glass transition temperature 233 – tackifier 233 copolymer – diblock 47, 229 – polycaprolactone 47 – radial 229 – statistical 421 – styrene-butadiene 389 – styrene-isoprene block 230 – tetrablock 229 – triblock 229 copper (Cu) 72, 89, 103, 137, 138, 445, 466 copper ions 84, 470 coupling agent 17 – silane 19 covalent bond 2 Cr 89 curing process 72 curing reaction – IR spectra 77 – polyurethane (PU) 77

c carbodiimide 84 carbon fiber reinforced epoxy 541 cell – adhesion 157 – endothelial 157 – fibroblast 157 chelate 89 – crystals 94 chemical degradation 479

d DDA s. dicyandiamide DDS s. diaminodiphenylsulfone dendrimers 198 density functional theory (DFT) 18, 20 – IR spectrum 75 density functional tight binding (DFTB) 18, 20 DETA s. diethylenetriamine

Subject Index DFT s. density functional theory DFTB s. density functional tight binding DGEBA s. diglycidyl ether of bisphenol A diaminodiphenylsulfone (DDS) 103 dicyandiamide (DDA) 479 Diels-Alder reaction – cycloaddition 289 – reaction, kinetics 289 diethylenetriamine (DETA) 90, 134, 445, 447, 466, 479 differential scanning calorimetry (DSC) 91, 254, 370, 483, 488, 489, 498, 499 diffusion 72 – compatibility, polymer 387 – contact temperature 383 – glass transition temperature 383, 387 – interdiffusion 387 – interfacial width 382 – kinetics 383 – miscibility, polymer 382 – mobility, polymer 382, 387 diglycidyl ether of bisphenol A (DGEBA) 4, 90, 103, 133, 219, 445, 447, 466, 479 diglycidyl ether of bisphenol A/diethylenetriamine (DGEBA/DETA) 126 diphenylmethane-4,4'-diisocyanate 72 dissipation, viscoelastic behavior 59 DSC s. differential scanning calorimetry durability 217, 479 – aging 445 – cathodic delamination 516 – chemical degradation 507 – corrosive degradation 516 – electrochemical/corrosive reactions 507

e EDX s. energy dispersive analysis of X-rays E-glass 89 eigenvibration 75 elastic bonding 525 elastomers 59, 387 energy-dispersive analysis of X-rays (EDX) 103 environmental scanning electron microscopy (ESEM) 163 epoxides – addition of alcohols 211 – cationic polymerization 205 – initiators 207 – photopolymerization 207 – polymerization kinetics 209 – traces of moistures 207

epoxy 205, 217, 407, 408, 409, 410, 412, 541, 555 – aging 445, 479, 492, 498, 500 – composite systems 221 – degradation 479, 491, 500, 501 – glass transition 498, 500 – – temperature 94 – hot-curing 495 – infrared spectroscopy 495 – network formation 94 – phase separation 479, 498, 500 – RT-curing 484 – stiffness 125 – toughened 196 – water uptake 500 epoxy-amine 89, 126, 516 – durability 89 – hypersonic acoustic properties 126 – mechanical properties 89 – polymer 105 – – glass transition temperature 100 – – l-IR spectroscopy 97 – – mechanical properties 95 – – metal ion diffusion 101 – – organometallic complexes 101 – – stoichiometric ratio 100 epoxy films 449 – aging 487, 490, 494, 498, 500 – cathodic delamination 516 – corrosive degradation 516 – curing 485 – degradation 491, 500, 501 – density 469 – free volume void 469 – free volume void size 468, 469 – glass transition 485, 494, 498, 500 – hot-curing 495 – hygro-thermal 469 – hygro-thermo-oxidative aging 450, 452 – infrared spectroscopy 495 – oxirane consumption on metals 448 – phase separation 498, 500 – positron annihilation lifetime spectroscopy (PALS) 468 – thermal aging 450, 452, 469 epoxy interphases, – epoxy 103 – mechanics 103 – modulus 103 – stiffness 103 epoxy-metal, dismantlability 559 epoxy resin 1

571

572

Subject Index ERAS s. external reflection absorption spectroscopy ESEM s. environmental scanning electron microscopy external reflection absorption spectroscopy (ERAS) 74

f failure surface 223 fatigue – test 365, 374 – interphase 365 Fe 89 fibers – Kevlar 217 – Zylon 217 force-distance curve 61 fouling, bioadhesion 176 fracture toughness 196 friction – dissipation mechanism 59 – interfacial interaction 59 – macroscopic 59 – nanoscopic 59 – polymers 59

g glass 59, 435 glass fiber reinforced epoxy 541 glass transition temperature 484, 543

h HBP s. hyperbranched polymers HMPSA s. hot-melt pressure-sensitive adhesive hot-melt pressure-sensitive adhesive (HMPSA) 229 – complex shear modulus 231 – formulation 229 hyperbranched polymers (HBP) 198, 217 – epoxies 217 – polyurethane 217

i ICP s. inductively coupled plasma spectroscopy immobilization 72 inductively coupled plasma spectroscopy (ICP) 92 infrared reflection-absorption spectroscopy (PM-IRRAS) 47 infrared spectroscopy 72, 271 – aging 445

– attenuated total reflection spectroscopy (ATR) 74, 163, 399, 483 – density functional theory 75 – eigenvibration 75 – external reflection absorption spectroscopy (ERAS) 74 – l-FTIR 91 – FTIR-ATR spectroscopy 249 – FT-NIR 92 – IR band assignment 75 – IR external reflection absorption spectroscopy (IR-ERAS) 74, 447, 483 – IR external reflection spectroscopy (IR-ERAS) 465 – IR microscopy 483 – IR reflection absorption spectroscopy (IRRAS) 93, 291 – IR spectra calculation 74 – IRRAS s. IR reflection absorption spectroscopy – micro-infrared spectroscopy (l-FTIR) 91 – near-infrared spectroscopy (FT-NIR) 92 – quantitative FTIR results 75 – transition dipole moment 75 interface – buried 8 – chemical degradation 507 – chemistry 1, 292 – electrochemical/corrosive reactions 507 – reaction 4, 292 interfacial bonding 2 – Fe 11 interlaminar shear strength 221 interphase 2, 3, 12, 71, 93, 395 – continuum mechanics 319 – crack 365 – crosslinking 387 – Cu/Epoxy 116 – diffusion 387 – – ions 514 – – water 514 – elastic properties 135 – electrochemistry 514 – energy release rate 381 – epoxy 125 – epoxy-Al 138 – epoxy-amine 89 – epoxy-Au 138 – epoxy-Cu 138 – epoxy-metal 137 – – internal stresses 141 – – stiffness 141 – epoxy-Mg 138

Subject Index – epoxy/silicone rubber 135 – fatigue 365 – fracture 365 – indentation techniques 110 – interdiffusion 399 – mapping 110 – mechanical properties 125, 319, 365, 387 – miscibility 381 – modulus 397 – PEO-TMPC 365 – PVP/Epoxy 118 – sound velocity profile 135 – stiffness 125, 397 – vitrification 135 IPDA s. isophoronediamine IR external reflection absorption spectroscopy (IR-ERAS) 447, 483 IR spectra calculation, optical function 74 IR-ERAS s. IR external reflection absorption spectroscopy iron 11, 435, 516 isocyanate 73, 220 isophoronediamine (IPDA) 90

l lap shear strength 559 lap shear test 541, 559 – fatigue 536

mechanical testing, three point flexure test 91 Mg 89, 138 microcapsules 555 micro-indentation 397 – artifacts 110 – depth-sensing (DSI) 108 microtomy, ultra low-angle (ULAM) 12 molecular mobility 72

n nanoindentation, viscoelastic contribution 33 nanophase separation 196 nanosystems 189 NDT s. nondestructive testing network structure 72 Ni 89 NMR imaging 435 NMR spectroscopy 435 nondestructive testing (NDT) – adhesive joints 403 – nonlinear ultrasonic transmission 403 – scanning Kelvin probe 507 – unilateral NMR 435 nonlinear ultrasonic testing, epoxy-aluminum 403

o m macromolecules, biocompatible 151 magnesium 89, 138, 220 mechanical properties 391, 410 – adhesive 542 – adhesive joints 542 – aging 542 – crack 337 – creep 545 – cyclic loading 542 – E-modulus 543 – fatigue 547 – interphase 125, 319 – nonlinear ultrasonic transmission 403 – Ogden model 544 – plastic deformation 543 – pressure-sensitive adhesive (PSA) 337 – quasi-static-loading 542 – static long-term loading 542 – stiffness 125 – strain hardening 543 – strength 543 – stress-strain curve 545 mechanical spectroscopy 231

optical microscopy, polarized (POM) 92 organometallic complex 89 organosilane 1 – c-glycidoxypropyltrimethoxysilane 5 oxide stripping 11 oxides 17

p PALS s. positron annihilation lifetime spectroscopy PAMAM s. poly(amido-amine) PCL s. polycaprolactone PDMS s. polydimethylsiloxane peel strength 219, 249, 265, 391 peel test 337 PET s. poly(ethylene terephthalate) phase separation 72 physical vapor deposition (PVD) 73 plasma modification 265 – pulsed plasma-initiated homo- or copolymerization 265 plasma polymer layers 265 plasma treatment 178, 265 – polymer surfaces 267

573

574

Subject Index – pulsed plasma polymerization 267, 289 – surface functionalization 267, 289 PM-IRRAS s. infrared reflection-absorption spectroscopy PMMA s. poly(methyl methacrylate) PnBA s. poly(n-butyl acrylate) polyaddition, polyurethane 72 poly(amido-amine) (PAMAM) 217, 219 poly(n-butyl acrylate) (PnBA) 421 polycaprolactone (PCL) 47 polydimethylsiloxane (PDMS) 34, 59, 60, 146 polyetherimide (PEI) 217 polyethylene (PE) 265 poly(ethylene oxide) (PEO) 365 poly(ethylene terephthalate) (PET) 161, 253 polyisoprene rubber 389 polymer films 47, 71, 365 – aging 445, 465, 479 – anodic delamination 518 – cathodic delamination 518 – curing 445, 479 – degradation 479 – epoxy 445 – epoxy-metal 465, 479 – free volume 465 – positron annihilation lifetime spectroscopy (PALS) 465 polymer surfaces – functionalization 267 – plasma treatment 267 – pulsed plasma polymerization 267 poly(methyl methacrylate) (PMMA) 5, 47, 179 poly(4-methyl-1-pentene) 133 polyol 76, 220 poly(p-phenylene-2,6-benzobisoxazole) 221 polypropylene (PP) 180, 196, 265 poly(propylene ether) – diol 73 – triol 73 polystyrene (PS) 369 polytetrafluoroethylene (PTFE) 159, 265 – biochemical modification 157 poly(tetramethyl bisphenol A polycarbonate) (TMPC) 365 polyurethane (PU) 71, 72, 398, 435, 541 – crosslinking 77 – cured films, chemical structure 80 – curing reaction 73 – post-cured films 80 – thin film morphology 79

– urea-like species 82 poly(vinylidene fluoride) (PVdF) 13 polyvinylpyrrolidone (PVP) 103, 179 POM s. optical microscopy positron annihilation lifetime spectroscopy (PALS), free volume 465 pressure-sensitive adhesive (PSA) 249, 337 – acrylate 251, 421 – block copolymer 337 – cavities 346 – chemical composition 252, 337 – copolymer 421 – fibrillar structure 356 – glass transition temperature 259 – interfacial crack 356 – interfacial fracture 356 – mechanical properties 337 – Mooney-Rivlin model 353 – nano-bubbles 421 – on ethylene-propylene copolymer 356 – peel test 337 – slip-tube model 353 – stress-strain curve 337, 344 – uniaxial tension 353 – UV-crosslinking 249 probe tack 249 prostheses, cardiovascular 157 proteins 184 – adhesive 143 – barnacle 143 – Min3 157 PS s. polystyrene PSA s. pressure sensitive adhesive PTFE s. polytetrafluoroethylene PU s. polyurethane pull-off test 371 PVD s. physical vapor deposition PVdF s. poly(vinylidene fluoride) PVP s. polyvinylpyrrolidone

q quantum mechanical modeling 75

r reconstruction – a-quartz (0001) 21 – a-SiO2 (0001) 19 – – surface 18 rheology 229 – copolymer 236 – copolymer blends 236 – hot-melt pressure-sensitive adhesive 229 – model 236

Subject Index

s SAM s. self assembled monolayer sandwich components, assembly 525 SAXS, s. small angle X-ray scattering scanning force microscopy (SFM) 72 – cellular force spectroscopy 157 – force modulation microscopy (FMM) 103 scanning Kelvin probe 507, 509 – blister test 521 sectioning – chemical 12 – taper 12 self assembled monolayer (SAM) 34, 291 SFM s. scanning force microscopy shear adhesion failure temperature (SAFT) 249 shear strength 219 shear test 319 silane 221 – (6-aminohexyl)aminopropyltrimethoxysilane 34 – (2-carbomethoxy)ethyltrichlorosilane 34 – (3-glycidyloxypropyl)trimethoxysilane 23 – hexadecyltrichlorosilane 34 – modified 541 – 1H, 1H, 2H, 2H-perfluorodecylmethyldichlorosilane 34 – triethoxyvinylsilane 23 silica, silicon oxide 18 silicon 33 silicon oxide – a-SiO2 (0001) surface 18 – silica 18 silsesquioxanes, polyhedral oligomeric (POSS) 200 simulation – atomistic 17 – continuum mechanics 319 – density functional theory (DFT) 18 – molecular mechanical 18 – quantum mechanical 19 – tight binding 18 single lap shear 219, 308 small angle X-ray scattering (SAXS) 342, 421, 423, 424 – during tack test 424 Sn 89 spin-coating 73 sputter depth profiling 11 steel – mild 11

– stainless 541 sterilization 178 stiffness mapping, scanning force microscopy (SFM) 103 styrene-butadiene copolymer 389 surface energy 33 surface functionalization 265, 289 surface treatment – CLP 541 – composite materials 305 – flame 541 – – pyrolysis, silane 541 – functionalization 265, 289 – in the production line 540 – laser 305, 541 – plasma 265, 289

t tack 230, 258, 374 tack test 62, 421, 424 tackifier 231, 343 TDI s. toluene diisocyanate TEM 10 tensile test 365, 410 – uniaxial 319 thermodynamic work of adhesion 33 three point bending 221 Ti 89 ToF-SIMS 1, 2, 6 toluene diisocyanate (TDI), urone 1, 6 T-Peel 219

u ultra-low-angle microtomy (ULAM) 12 urea-like species 82 uretdion 84 urethoneimine triisocyanate 72

w wedge test 225

x XPS s. X-ray photoelectron spectroscopy X-ray diffraction (XRD) 92 X-ray photoelectron spectroscopy (XPS) 1, 181, 271, 465, 541 – labeling of functional groups 272 XRD s. X-ray diffraction

z Zn 89

575

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