Spectroscopy of Rubber and Rubbery Materials

Spectroscopy of Rubbers and Rubbery Materials Victor M. Litvinov and Prajna P. De Spectroscopy of Rubbers and Rubbery ...

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Spectroscopy of Rubbers and Rubbery Materials Victor M. Litvinov and Prajna P. De

Spectroscopy of Rubbers and Rubbery Materials

Victor M. Litvinov and Prajna P. De

Rapra Technology Limited Shawbury, Shrewsbury, Shropshire, SY4 4NR, United Kingdom Telephone: +44 (0)1939 250383 Fax: +44 (0)1939 251118 http://www.rapra.net

First Published 2002 by

Rapra Technology Limited Shawbury, Shrewsbury, Shropshire, SY4 4NR, UK

©2002, Rapra Technology Limited

All rights reserved. Except as permitted under current legislation no part of this publication may be photocopied, reproduced or distributed in any form or by any means or stored in a database or retrieval system, without the prior permission from the copyright holder. A catalogue record for this book is available from the British Library.

Every effort has been made to contact copyright holders of any material reproduced within the text and the authors and publishers apologise if any have been overlooked.

ISBN: 1-85957-280-4

Typeset by Rapra Technology Limited Printed and bound by Polestar Scientifica, Exeter, UK

Contents

Contents 1

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric Analysis Techniques ......................................................... 1 1.1

Introduction ........................................................................................... 1

1.2

Instrumental Part ................................................................................... 3 1.2.1

1.3

2

Experimental ........................................................................... 12

Applications of Hyphenated Thermogravimetric Analysis Techniques ........................................................................................... 14 1.3.1

Thermogravimetry - Fourier Transform Infrared Spectroscopy ............................................................................ 19

1.3.2

Thermogravimetry-Mass Spectrometry .................................... 25

1.3.3

Thermogravimetry - Differential Scanning Calorimetry ........... 29

1.3.4

Thermogravimetry - Differential Scanning Calorimetry - Mass Spectrometry ................................................................ 30

1.3.5

Thermogravimetry - Differential Thermal Analysis ................. 30

1.3.6

Thermogravimetry - Differential Thermal Analysis - Mass Spectrometry ................................................................ 33

1.3.7

Thermogravimetry - Gas Chromatography - Mass Spectrometry ................................................................ 34

1.3.8

Thermal Desorption - Gas Chromatography - Fourier Transform Infrared Spectroscopy/Mass Spectrometry .............. 34

1.4

Future Prospects of Hyphenated Thermogravimetric Techniques in Elastomer Characterisation .............................................................. 35

1.5

Summary .............................................................................................. 36

Photoacoustic Fourier Transform Infrared Spectroscopy of Rubbers and Related Materials ................................................................................... 49 2.1

Introduction ......................................................................................... 49

2.2

History of Photoacoustic Spectroscopy ................................................ 49

i


3

4

ii

2.3

Theory of Photoacoustic Spectroscopy ................................................ 50

2.4

Instrumentation for PA-FTIR Analyses ................................................ 52

2.5

Analysis of Carbon-Filled Rubbers ...................................................... 57

2.6

Quantitative Analysis of Polymers ....................................................... 60

2.7

Surface Analysis and Depth Profiling ................................................... 66

2.8

Determination of Orientation Function of Polymeric Materials ........... 71

2.9

Conclusion ........................................................................................... 71

Infrared Spectroscopy of Rubbers ................................................................. 77 3.1

Introduction ......................................................................................... 77

3.2

Sample Preparation .............................................................................. 80

3.3

Different Types of IR Spectroscopy ...................................................... 81

3.4

Quantitative Analysis ........................................................................... 81

3.5

Applications of IR Spectroscopy .......................................................... 82 3.5.1

Rubber Blends ......................................................................... 92

3.5.2

Self-crosslinking Blends............................................................ 96

3.5.3

Polyurethanes .......................................................................... 98

3.5.4

Rubber-filler Interaction ........................................................ 102

3.5.5

Milling ................................................................................... 108

3.5.6

Adhesion ................................................................................ 111

3.5.7

Degradation ........................................................................... 113

3.6

Reverse Engineering ........................................................................... 113

3.7

Conclusion ......................................................................................... 114

Application of Infrared Spectroscopy to Characterise Chemically Modified Rubbers and Rubbery Materials .................................................. 125 4.1

Introduction ....................................................................................... 125

4.2

The Infrared Spectra of Commonly Used Diene Rubbers ................... 125

Contents

4.3

Hydrogenation ................................................................................... 126

4.4

Halogenation ..................................................................................... 135

4.5

Isomerisation ..................................................................................... 139

4.6

Cyclisation ......................................................................................... 139

4.7

Hydrosilylation .................................................................................. 142

4.8

Hydroboration ................................................................................... 144

4.9

Hydroformylation .............................................................................. 144

4.10 Oxidation .......................................................................................... 146 4.11 Phosphonylation ................................................................................ 146 4.12 Sulfonation ........................................................................................ 147 4.13 Ionomer Formation ............................................................................ 147 4.14 Ionomeric Blends ............................................................................... 151 4.15 Weathering and Degradation of Polymers .......................................... 153 4.16 Modification by Radiation ................................................................. 156 4.17 Conclusion ......................................................................................... 157 5

6

Infrared Spectroscopy of Rubbery Materials ............................................... 167 5.1

Introduction ....................................................................................... 167

5.2

Polyethylenes ..................................................................................... 168

5.3

Polyvinyl Chloride ............................................................................. 180

5.4

Thermoplastic Elastomers .................................................................. 188

5.5

Conclusion ......................................................................................... 200

Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy ............................................................................................... 207 6.1

Introduction ....................................................................................... 207 6.1.1

General Introduction to EPDM ............................................. 207

iii


6.2

6.1.2

EPDM Crosslinking ............................................................... 208

6.1.3

Studies into the Chemistry of Rubber Crosslinking ............... 209

6.1.4

Scope ..................................................................................... 210

Sulfur Vulcanisation ........................................................................... 211 6.2.1

6.3

6.4 7

Peroxide-curing .................................................................................. 225 6.3.1

General .................................................................................. 225

6.3.2

Polydiene Elastomers ............................................................. 226

6.3.3

EPDM .................................................................................... 227

Concluding Remarks and Future Outlook ......................................... 237

NMR Imaging of Elastomers ...................................................................... 247 7.1

7.2

7.3

7.4

iv

Sulfur Vulcanisation of Polydiene Rubbers ............................ 211

NMR Imaging and Contrast .............................................................. 248 7.1.1

Principle of Fourier NMR ...................................................... 248

7.1.2

Spatial Resolution .................................................................. 251

7.1.3

Contrast ................................................................................. 253

Applications ....................................................................................... 264 7.2.1

Defects and Heterogeneities in Technical Elastomer Products ................................................................................. 265

7.2.2

Covulcanisation ..................................................................... 269

7.2.3

Blending ................................................................................. 271

7.2.4

Crosslink Density ................................................................... 271

7.2.5

Vulcanisation Process ............................................................. 273

7.2.6

Ageing ................................................................................... 274

7.2.7

Sample Deformation .............................................................. 275

Spatially Resolved NMR .................................................................... 276 7.3.1

The NMR-MOUSE ................................................................ 277

7.3.2

Applications ........................................................................... 279

7.3.3

Imaging with the NMR-MOUSE ........................................... 282

Summary ............................................................................................ 282

Contents

8

NMR in Soft Polymeric Matter: Nanometer Scale Probe ............................ 291 8.1

Introduction ....................................................................................... 291

8.2

Polymeric Networks .......................................................................... 292

8.3

8.4

8.5

8.6

8.2.1

Molten High Polymers ........................................................... 292

8.2.2

Crosslinked Chains ................................................................ 292

8.2.3

Semi-crystalline Polymers....................................................... 292

8.2.4

Block Copolymers .................................................................. 293

8.2.5

Loaded Polymers ................................................................... 293

8.2.6

Aggregated Polymers ............................................................. 294

8.2.7

Network Distribution Function ............................................. 294

Basis of the NMR Approach .............................................................. 295 8.3.1

Chain Elongation ................................................................... 296

8.3.2

NMR Evidence for Networks: Pseudo-solid Spin-echoes ....... 298

Crosslinked Chains ............................................................................ 299 8.4.1

End-linked Calibrated Chains ................................................ 299

8.4.2

Characteristic NMR Rates ..................................................... 299

8.4.3

Strand Length Dependence .................................................... 301

8.4.4

Randomly Crosslinked Chains ............................................... 302

8.4.5

Latex Suspensions .................................................................. 306

8.4.6

Kinetics of Gelation ............................................................... 307

Polymeric Crystallisation ................................................................... 307 8.5.1

Crystallisation-NMR Relationship ........................................ 307

8.5.2

Kinetics of Crystallisation ...................................................... 308

Entangled High Polymers ................................................................... 309 8.6.1

Temporary Networks ............................................................. 310

8.6.2

Short Chain Dynamics ........................................................... 312

8.6.3

Long Chain Dynamics ........................................................... 313

8.7

Adsorption on Mineral Aggregates .................................................... 315

8.8

Conclusion ......................................................................................... 317

v


9

Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR ......................................................................................... 321 9.1

Introduction ....................................................................................... 321

9.2

Sulfur Vulcanisation Mechanism ........................................................ 322

9.3

The NMR Methods for Assigning Resonances to Chemical Structure ............................................................................ 324

9.4

Unaccelerated Sulfur-vulcanisation of NR ......................................... 327

9.5

Accelerated Sulfur-vulcanisation of NR and IR .................................. 327

9.6

Sulfur-vulcanisation of BR ................................................................. 333

9.7

Sulfur-vulcanisation of SBR ............................................................... 338

9.8

Peroxide, Radiation, and High Pressure Vulcanisation ...................... 338

9.9

Vulcanisation of Other Elastomer Systems ......................................... 340

9.10 Effect of Carbon Black on Vulcanisation of Elastomers ..................... 341 9.11 Effect of Silica Filler on Vulcanisation Chemistry .............................. 347 9.12 Thermal-Oxidation of Network Structures ........................................ 347 9.13 Summary ............................................................................................ 348 10 Characterisation of Chemical and Physical Networks in Rubbery Materials Using Proton NMR Magnetisation Relaxation ........................... 353 10.1 Introduction ....................................................................................... 353 10.2 Network Structure Analysis by Means of NMR Transverse Magnetisation Relaxation ................................................ 355 10.3 Characterisation of Network Heterogeneity and Network Defects ................................................................................ 360 10.4 Network Structure in Oil-Extended Rubbers - Effect of Chain Entanglements......................................................... 366 10.5 Network Structure in Filled Rubbers - Rubber-Filler Interface and the Structure of the Physical Network .......................... 368

vi

Contents

10.5.1 NMR Relaxation of Filled Rubbers ....................................... 368 10.5.2 Carbon-Black-Filled Rubbers ................................................. 369 10.5.3 Silica-Filled Silicon Rubbers ................................................... 374 10.5.4 Silica-Filled Conventional Rubbers ........................................ 378 10.6 Chains Grafted onto a Filler Surface .................................................. 379 10.7 Semi-Crystalline Elastomers ............................................................... 381 10.8 Ionic Viscoelastic Materials ............................................................... 383 10.9 Rubbery Phases in Blends and Emulsions .......................................... 384 10.10Real-Time NMR Experiments............................................................ 385 10.11Low-Field NMR Magnetisation Relaxation Experiments for Quality Control Purposes ............................................................. 387 10.12Conclusions ....................................................................................... 388 11 High Resolution NMR of Elastomers ......................................................... 401 11.1 Introduction ....................................................................................... 401 11.2 Structural Feature of Elastomers ........................................................ 402 11.3 Analysis of Chemical Composition and Sequence Distribution .......... 404 11.3.1 Accuracy of NMR Measurements of Chemical Composition .......................................................................... 404 11.3.2 Analysis of Chemical Composition Distribution using SEC-NMR .................................................................... 410 11.3.3 Analysis of Sequence Distribution .......................................... 413 11.3.4 Analysis of Chemically Modified Structure and Graft Polymers ................................................................ 419 11.4 Analysis of End-groups and Branching .............................................. 422 11.4.1 Assignment of Small Signals .................................................. 422 11.4.2 Functionality of Telechelic Diene Polymers ............................ 423 11.4.3 Structure of Terminal Groups ................................................ 424 11.4.4 Structure of Branch-points and Coupling Points .................... 426

vii


11.5 Structural of Naturally Occurring Polyisoprenes ............................... 429 11.5.1 Structure of Natural cis- and trans-Polyisoprenes .................. 429 11.5.2 Structure of Natural Rubber .................................................. 433 11.6 Application of High-Resolution NMR ............................................... 436 11.6.1 Multinuclear High-resolution NMR ...................................... 436 11.6.2 NMR of Swollen State and Latex .......................................... 443 11.7 Conclusion ......................................................................................... 448 12

129

Xe NMR of Elastomers in Blends and Composites.................................. 457

12.1 Introduction to 129Xe NMR Spectroscopy of Materials ...................... 457 12.1.1

129

Xe NMR Spectroscopy ....................................................... 457

12.1.2 NMR of 129Xe in the Gas Phase ............................................. 458 12.1.3

129

Xe NMR of Polymers ......................................................... 459

12.1.4

129

Xe Pulsed Field Gradient Echo (PFGE) Spectroscopy ......... 462

12.2 Experimental ..................................................................................... 463 12.3

129

12.4

129

Xe NMR of Ethylene/Propylene Copolymers in Blends and Block-copolymers with Polypropylene ........................................ 465 Xe NMR of an Ethylene/Propylene/Diene Terpolymer in Carbon Black Composites .............................................................. 474

12.5 The Self-diffusion Coefficient of Xe in Elastomers ............................. 480 12.5.1 General .................................................................................. 480 12.5.2 Xe Diffusion in a iPP/EPDM Blend ........................................ 482 12.5.3 Xe Diffusion as a Function of the Diffusion Time .................. 485 12.6 Conclusions ....................................................................................... 487 13 Swollen Rubbery Materials: Chemistry and Physical Properties Studied by NMR Techniques ...................................................................... 491 13.1 Introduction - The Swelling of Crosslinked Rubbers ......................... 491 13.1.1 The Theory of Rubber Swelling ............................................. 491 13.1.2 Relationship with NMR Parameters ...................................... 492

viii

Contents

13.1.3 Reviews of NMR of Crosslinked Rubbers and Polymers ......................................................................... 493 13.2 Motion of Polymer Chains in Polymer Solutions and Swollen Networks .............................................................................. 494 13.2.1 Background ........................................................................... 494 13.2.2 Solutions and Bulk Rubbers ................................................... 494 13.3 Motion of Small Molecules in Swollen Rubbers ................................ 496 13.3.1 Self-Diffusion of Small Molecules in Rubbers ........................ 496 13.3.2 Self-Diffusion of Rubbers ...................................................... 498 13.4 NMR Imaging of Swollen Rubbers .................................................... 499 13.4.1 Macroscopic Diffusion of Small Molecules in Swollen Rubbers ................................................................ 499 13.4.2 NMR Imaging of Swollen BR and Polyisoprene Rubbers ...... 500 13.5 Studies of Network Density in Swollen Rubbers and Blends .............. 502 13.5.1 Measurements of Transverse Relaxation Times ..................... 502 13.5.2 Estimation of Crosslink Density from NMR Linewidths ....... 506 13.5.3 High-Resolution 13C MAS NMR of Rubbers ......................... 510 13.6 Summary ............................................................................................ 511 14 Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials ............................................................................. 519 14.1 Introduction ....................................................................................... 519 14.2 Basics of NMR in Viscoelastic Polymers ............................................ 520 14.2.1 Anisotropic Spin Interactions ................................................. 520 14.2.2 Manipulation of Spin Interactions ......................................... 522 14.2.3 Residual Couplings and Dynamic Order Parameters ............. 524 14.2.4 One-dimensional NMR Studies of Molecular Motions and Dynamic Order ................................................. 526 14.3 Multidimensional NMR Spectroscopy of Viscoelastic Materials........ 529 14.3.1 Principle of Multidimensional NMR ..................................... 529

ix


14.3.2 Two-dimensional NMR Techniques and Applications to Viscoelastic Polymers .................................... 530 14.5 Conclusion .......................................................................................... 550 15 Deuterium NMR in Rubbery Materials ...................................................... 557 15.1 Introduction ....................................................................................... 557 15.2 2H NMR Background ........................................................................ 559 15.3 Model, Labelled Rubber Systems ....................................................... 561 15.3.1 General Presentation: End-linked versus Randomly Crosslinked Networks ........................................................... 561 15.3.2 Pseudosolid Behaviour ........................................................... 561 15.3.3 Dangling Chains .................................................................... 563 15.3.4 Time Scale of Motions ........................................................... 563 15.3.5 Uniaxially Deformed Model Networks .................................. 566 15.3.6 Physical Origin of the Induced Orientation............................ 573 15.3.7 Correlation to Elastic Properties ............................................ 575 15.4 Using Deuterated Probes .................................................................... 579 15.5 Filled, Composite and Thermoplastic Elastomers .............................. 582 15.5.1 Filled Elastomers .................................................................... 582 15.5.2 Semi-crystalline Polymers....................................................... 584 15.5.3 Thermoplastic Elastomers ...................................................... 585 15.6 Conclusion .......................................................................................... 588 Abbreviations .................................................................................................... 597 Contributors ...................................................................................................... 607 Index ................................................................................................................. 611

x

Preface

Preface

This book deals with the application of spectroscopic techniques for characterisation of chemical and physical structures in viscoelastic materials, such as unvulcanised elastomers and their vulcanisates, various rubbery materials and some plastics, which when blended with particular additives (plasticisers) behave like rubbers. The common feature of all these materials is a low glass transition temperature and a molecular network, which is formed due to chemical and/or physical junctions between long chain molecules. Depending on the applied force this network, neither liquid nor solid, reveals both solidlike and liquid-like properties. This dual behaviour determines the application areas of these materials. A distinguishing feature of rubbers is good elasticity, which is largely determined by the network structure and phase/components composition of the material. Despite the apparent simplicity, the network structure has a complex topology, which can significantly affect the functional properties. Different types of network heterogeneity are formed in the cured material, if no special precautions are taken to control the curing chemistry and conditions. A difference in curing conditions through the sample volume, such as the content of vulcanisation agents and temperature, results in spatial heterogeneity of the network structure. Ageing and weathering can be other reasons for the spatial heterogeneity. The structure of initial rubber and vulcanisation chemistry could cause molecular scale heterogeneity of the network. Analysis of the rubbery materials is complicated by the fact that rubbery products, such as tyres, tubes, seals, V-belts and hoses, contain in the rubbery matrix a significant amount of various compounds, i.e., fillers, vulcanising agents, antioxidants and plasticisers. Due to the complex composition, no single technique can provide a good understanding of the effect of chemical and physical structures on the functional properties of rubbery materials. It is acknowledged that traditional methods, such as mechanical measurements and equilibrium swelling, are not capable of giving reliable information about the network structure and phase composition of rubbery materials, since these methods determine volume average material properties. Local methods, which probe molecular properties, are well suited for this purpose. Apparently, the most comprehensive information on chemical and physical structures in the relation to material properties can be obtained by using a combination of macroscopic techniques and methods that provide information on the molecular level.

xi

Spectroscopy of Rubbers and Rubbery Materials Spectroscopy has become a powerful tool for the determination of polymer structures. The major part of the book is devoted to techniques that are the most frequently used for analysis of rubbery materials, i.e., various methods of nuclear magnetic resonance (NMR) and optical spectroscopy. One chapter is devoted to (multi) hyphenated thermograviometric analysis (TGA) techniques, i.e., TGA combined with Fourier transform infrared spectroscopy (FT-IR), mass spectroscopy, gas chromatography, differential scanning calorimetry and differential thermal analysis. There are already many excellent textbooks on the basic principles of these methods. Therefore, the main objective of the present book is to discuss a wide range of applications of the spectroscopic techniques for the analysis of rubbery materials. The contents of this book are of interest to chemists, physicists, material scientists and technologists who seek a better understanding of rubbery materials. The book brings together the various spectroscopic techniques for obtaining the following information: Chemical structure of rubbery materials: Chapters 1, 2, 3, 4, 5, 6, 9 and 11, describe applications of (multi) hyphenated TGA techniques, optical and high-resolution NMR spectroscopes for the analysis of chain microstructures and conformations, chemical composition of components, additives and volatiles in rubbery materials, vulcanisation chemistry, functional groups analysis and chemical modification of rubbery materials. Network structure analysis is discussed in Chapters 7, 8, 10 and 13. These chapters deal with the characterisation of the structure of chemical and physical networks, rubberfiller physical network, network defects and its heterogeneity using NMR relaxation techniques and NMR imaging. Heterogeneity of rubbery materials: Use of NMR imaging, NMR MOUSE (mobile universal surface explorer) and photoacoustic FT-IR is reviewed in Chapters 2 and 7. These methods allow determination of defects and heterogeneity in technical elastomer products, surface and depth profiles. Physical properties of rubbery materials are discussed in Chapters 5, 7, 8, 10, 12, 13, 14 and 15, which cover the following topics: phase structure and composition, molecular mobility, strain induced phenomena, phase separation and interfaces in rubbery materials. Functional properties and stability of rubbery materials: Chapters 1, 3, 4, 7, 12 and 13, give examples of applications of spectroscopic techniques for the characterisation of thermal stability and degradation, kinetics of thermal decomposition, ageing, oxidation and weathering, self-diffusion of small molecules in rubbery materials, adhesion of rubbers to metals, fluid adsorption and swelling. xii

Preface Processing of rubbery materials: The effect of milling on the heterogeneity of rubbery materials and changes in their chemical and physical structures is determined by means of optical spectroscopy, NMR relaxation experiments and NMR imaging (Chapters 3, 7 and 10). Welding of crosslinked polyethylenes is discussed in Chapter 5. Quality control: Applications of NMR relaxation experiments, NMR imaging, NMR MOUSE and (multi) hyphenated TGA techniques for quality control are discussed in Chapters 1, 7 and 10. As with any work of this type, this book is a result of a collective effort of many people. The editors thank all the authors for their excellent contributions. We are indebted to Ms. Frances Powers, Commissioning Editor of Rapra Technology, for her support in the course of the work during last three years. Many thanks go to Ms. Claire Griffiths and Ms. Sandra Hall for transforming the manuscripts into book format, and to all members of RAPRA editorial group. We would also like to thank Ms. Caroline Barlow who prepared the Index. Lastly, thanks to the family members of the editors, without whose help and understanding the book could not come to completion. Victor M. Litvinov and Prajna P. De 2002

xiii


xiv

1

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric Analysis Techniques Jan C.J. Bart and Charles Raemaekers

1.1 Introduction Thermal analysis is a group of techniques in which a physical property of a substance is measured as a function of temperature when the sample is subjected to a controlled temperature program. Single techniques, such as thermogravimetry (TG), differential scanning calorimetry (DSC), dynamic mechanical analysis (DMA), dielectric thermal analysis, etc., provide important information on the thermal behaviour of materials. However, for polymer characterisation, for instance in case of degradation, further analysis is required, particularly because all of the techniques listed above mainly describe materials only from a physical point of view. A hyphenated thermal analyser is a powerful tool to yield the much-needed additional chemical information. In this paper we will concentrate on simultaneous thermogravimetric techniques. Thermogravimetry measures total mass changes under precise heating conditions. Thermal events bringing about a change in the mass of a sample are adsorption and absorption, desorption (diffusion), outgassing, dehydration or desolvation, foaming, vapourisation, sublimation, decomposition, oxidation, reduction and other chemical reactions. These phenomena can therefore be studied by means of TG. Some factors influencing component loss are: chemical stability, volatility, rate of diffusion, sample thickness, flow rate, temperature, heating rate, etc. Apart from sample weight changes no other chemical or physical information is gathered about the matrix. For example, phenomena such as polymerisation and cross-linking cannot be detected by TG. Although thermogravimetric analysis is a powerful tool for quantitative polymer analysis its scope is further limited as also no information is obtained about the qualitative aspects of the evolved gases during polymer degradation. Both simultaneous and sequential techniques have been developed for the purpose of identification of these gases and volatiles. The great advantage of simultaneous thermoanalytical methods is that they allow the results to be interpreted and evaluated jointly. As may be seen from Figure 1.1 [1-27] hyphenated and multihyphenated methods greatly enlarge the scope of TG and are becoming important analytical tools for materials

1


Figure 1.1 (Multi) hyphenated thermogravimetric analysis techniques. The figures refer to references

characterisation. Some thirty different experimental approaches may already be identified, many with a proven record. The present trend towards ever increasing product quality standards demands greater sensitivity, accuracy, reproducibility and therefore constitutes a challenge to thermogravimetric analytical techniques. The thermal characterisation of elastomers has recently been reviewed by Sircar [28] from which it appears that DSC followed by TG/DTG are the most popular thermal analysis techniques for elastomer applications. The TG/differential thermal gravimetry (DTG) method remains the method of choice for compositional analysis of uncured and cured elastomer compounds. Sircar’s comprehensive review [28] was based on single thermal methods (TG, DSC, differential thermal analysis (DTA), thermomechanical analysis (TMA), DMA) and excluded combined (TG-DSC, TG-DTA) and simultaneous (TG-fourier transform infrared (TG-FTIR), TG-mass spectroscopy (TG-MS)) techniques. In this chapter the emphasis is on those multiple and hyphenated thermogravimetric analysis techniques which have had an impact on the characterisation of elastomers. The review is based mainly on Chemical Abstracts records corresponding to the keywords elastomers, thermogravimetry, differential scanning calorimetry, differential thermal analysis, infrared and mass spectrometry over the period 1979-1999. Table 1.1 contains the references to the various combined techniques.

2

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric…

Table 1.1 Selected references to (multi) hyphenated thermogravimetric analyses of elastomers Technique

Reference number

TG-FTIR

[29-48]

TG-MS

[25, 48-54]

TG-DSC

[55-62]

TG-DSC-MS

[50, 63]

TG-DTA

[57, 64-84, 149]

TG-DTA-FTIR

[39]

TG-DTA-MS

[59-62, 85]

TG-Tenax-TD-GC-FTIR-MS

[86]

1.2 Instrumental Part Thermogravimetric measurements require a thermobalance. There are many different types of TG analysers, varying in design (furnace, atmosphere, pressure, temperature range, sample quantity, sensitivity, resolution, etc.), degree of microcomputer control of the hardware, capabilities of the software, etc., all with typical specifications. The basic TG experiment consists of recording the weight of a sample as it is heated in a given environment either isothermally or at a controlled heating rate. The experimental record is a plot (thermal curve) of some form of weight change versus time or temperature of the sample. The simple additional step of using the derivative of the primary weight change (DTG) extends the capability and scope of the analysis. Heating of the sample may be combined with measurement of any parameter which further characterises the thermally activated transformation, such as heat flow (DTA, DSC), mass (DTG) or total gas flow (evolved gas analysis by MS or FTIR). The family of thermogravimetric techniques, as shown in Figure 1.1, can be divided into simultaneous and non-simultaneous methods. Non-simultaneous methods imply no realtime generation of thermogravimetric and other physico-chemical information. A continuous analysis of a gas stream from a thermal analyser is not readily possible by gas chromatography (GC); usually one has to use a gas-sampling loop for intermittent operation. Non-simultaneous methods usually involve cold traps (CT) or absorbent charcoal (Tenax). Because of the intermittent function of the analysis the coupling of GC to a thermogravimetric analyser is therefore still not so common. The separation function

3

Spectroscopy of Rubbers and Rubbery Materials by time in GC is, to a small extent, also given by thermogravimetry, as the gases are ‘temperature separated’ during analysis. Non-simultaneous TGA techniques, e.g., TGGC-MS, TG-Tenax-FTIR-MS, add surplus value and understanding when non-continuous data are sufficient. Due to the experimental set-up, these methods give inherently less insight in degradation behaviour than continuous methods. In particular, there is no continuous monitoring of evolved gases. In fact, only information about a specific point of the TG curve is collected at a given time. Discontinuous TG-GC, TG-GC-MS and TGTenax-TD-GC-MS are obviously inherently slower than TG-MS. The possibilities are improving though by using a fast column and an accurate fast sampling valve. However, only the fragments that are non-reactive, thermally stable, and volatile can be analysed by GC, whereas MS has none of these limitations. Despite the utility of TG-FTIR and TG-MS, a distinct disadvantage is that the presence of components at very low concentrations may be masked by higher concentration interferants. Thus, some researchers have incorporated the separation power of gas chromatography by collecting products in a trap or on the head of a capillary column for all or part of the TG run [87, 88]. However, these methods necessarily result in the loss of the time/temperature evolution data for the products analysed. McClennen and co-workers [20] have described TG-GC-MS and TG-GC-IR illustrating the feasibility of using a high-speed ‘transfer line’ GC technique for near real-time process monitoring applications using MS [89]. The approach is thus to transport products from the TG to the GC where they are rapidly separated on a short capillary column. Combination of a pulsed automated vapour sampling inlet and transfer line type GC column permits high-speed GC identification of individual TG products while maintaining sufficiently high temporal resolution with a ~ 1 minute sampling interval. By selecting columns to provide short retention times (< 1 minute) consistent with time-resolved profiles of the TG curve for comparison to the DTG curves the resulting TG-GC-MS and TG-GC-IR configurations are capable of separating thermal decomposition products, while still allowing characterisation of the evolution profile for each component. The classic chromatographic trade-off between efficiency and analytical time can be balanced in such a way to provide both real-time thermal evolution profiles of multiple components and separation sufficient to allow a significant degree of component identification by means of TG-GC-IR and TG-GC-MS. Meuzelaar and co-workers [90] have described an on-line high pressure TG-GC-MS system which requires only very small amounts (10-100 mg) of sample and can be operated at high pressure under different atmospheres (N2, He, H2, etc.). The system has been used in recycling lower grade post-consumer polymers, as coloured polyethylene and polystyrene or used rubber tyres, by co-processing with coal. The main characteristics of non-simultaneous methods, TG-GC and techniques requiring a cold trap, are given in Table 1.2.

4


Table 1.2 Characteristics of some non-simultaneous TG methods [26] TG-GC

TG-CT

• Roughly equivalent to PyGC plus weight change.

• Intermitted sampling allows for discontinuous analysis of evolved gases.

• Permits the analysis of components of different volatilities.

• Continuous monitoring of the weight change process and gas analysis is impossible.

• GC allows analysis only of those fragments that are non-reactive, thermally stable and volatile

• Recombination between products prior to subsequent analysis is not to be excluded.

It is especially important to have simultaneous techniques incorporating the most frequently used thermal methods. Early work allowed joint measurement of TG, DTG, and DTA. Paulik and co-workers [91, 92] devised simultaneous TG/DTG-DTA (the derivatograph). Later, also TG/DTG-DTA-MS has been described [93]. TG-DSC (or TGDTA) allows simultaneous measurements for the determination of the mass change (TG) and the energetic changes (DSC/DTA) on one sample under identical test conditions. As with this method all factors which influence the measurement signals, e.g., atmosphere, sample structure, temperature gradient, diffusion paths and packing density, are the same, TG and DSC/DTA results can be correlated and interpreted more easily. The complementary information obtained allows differentiation between endothermic and exothermic events which have no associated weight loss, e.g., melting and crystallisation, and those which involve a weight loss, e.g., degradation. Simultaneous TGA-DTA results are most often used to improve the interpretation of thermal events. Redfern [94] summarises the advantages of single sample simultaneous TG-DSC (or TGDTA) as follows: •

Greater efficiency (sample preparation time, run set-up time, instrument time).

•

Reduced influence of sample preparation (a single sectioning of the original material).

•

Higher accuracy of TGA temperature calibration (typically 0.1 °C for DSC and only 2 °C for stand-alone TGA).

•

Simplified interpretation for perturbation of results arising from thermal history, orientation effects, heat treatment, pressure during cutting, etc.

•

Combined evaluation assuring identical experimental and sampling conditions for both measurements with elimination of sources of uncertainty.

5

Spectroscopy of Rubbers and Rubbery Materials •

Correlation of observed effects.

•

Weight measurements validating quantitative measurements from DSC experiments.

•

Detection of moisture and determination of in-situ dry weight.

•

Quantitative measurements (based on correct dry sample mass and accurate enthalpic measurements).

The ability to obtain two thermal measurements on a single material is useful for rapidly comparing different batches of product or raw material in quality control. Simultaneous TG-DTA provides valuable information even in materials where no weight changes occur over the temperature range studied. A disadvantage of TG-DSC (or TG-DTA) is that the data obtained give no direct information on the nature of the chemical species involved. A typical modern assembly for simultaneous TG-DSC (or ‘STA’) is shown in Figure 1.2. Coupled instruments each need to operate under optimum conditions. Key elements in performance are the interface system and an integrated software package. Interface design is critical in allowing rapid mass transfer while not degrading the species of interest being transferred. No standardised thermogravimetric couplings are available. Descriptions of interfaces and interface techniques in coupled instrumentation have been reported: TG-FTIR [29, 30, 96], TG-MS [25, 49, 50, 63, 96], TG-DSC [55, 56], TG-DSCMS [50, 63], TG-DTA [56], TG-DTA-MS [93], TG-GC-MS [96], TD-GC-FTIR [97], TG-Tenax-TD-GC-FTIR-MS [86]. For identification purposes the classical techniques of infrared spectroscopy and MS are highly suited in hyphenation to thermogravimetry. Table 1.3 shows the main characteristics of TG-FTIR and TG-MS. Kaisersberger and co-workers [96] discussed hyphenation of a thermo-balance to infrared spectrometers. In this instrumental coupling, the infrared radiation must be brought into an intensive contact with the gas flow including the evolved gases, which is achieved in specially designed gas-measuring cells. FTIR spectrometers provide the high scan speeds which are required for the on-line coupling with thermal analysers for continuous monitoring of the gas composition. FTIR coupling is made by connection of the gas cell with the shortest possible heated transfer lines to the gas flow system of the thermobalance. As opposed to the coupling of mass spectrometers, the whole gas flow from the thermobalance should pass through the gas cell of the infrared spectrometer. While

6


Figure 1.2 Simultaneous TG-DSC Reproduced with kind permission of Netzsch-Gerätebau GmbH, Selb. Germany [95]

experimental results obtained from coupling of TG and FTIR spectroscopy were reported in 1981 [19], dedicated TG-FTIR instruments were not commercially available until 1987 [98-100]. Up to that time most work on the identification of evolved gases from TG had been in the TG-MS combination and reports on polymer studies using hyphenated TG-FTIR were relatively scarce [101, 102]. Several approaches to coupling TG and FTIR instruments have been reported [19, 29, 31, 55, 57, 103-108]. In conventional commercially available TG-FTIR systems, the evolved gases are led from the TG system to the spectrometer via a heated transfer line by a carrier gas flow [19, 103-105]. Another approach is the use of helium carrier gas at high flow rates, leading to the formation of an aerosol of the evolved components [31, 109]. The aerosol is then led to the spectrometer without loss of high-molecular-mass components. This system performs quantitative measurements and preserves and monitors very high molecular weight condensibles. In an ‘on-the-spot’ TG-FTIR technique the radiation is brought to the TG system - as opposed to bringing the evolved components to the spectrometer [29, 110]. The IR beam is led directly to the TG system, where it is reflected by a mirror mounted inside the TG equipment and is subsequently detected by the standard FTIR detector. Compared to

7


Table 1.3 Characteristics of TG-FTIR and TG-MS methods [26] TG-FTIR • Characterises and quantifies gaseous decomposition products • Permits specific compound analysis: fingerprint identification based on vibrational spectra • Identifies classes of compounds • Requires IR absorbency • Makes use of reference spectra • Is of limited sensitivity only ª Has lower compound specificity than TG-MS (sub-μg levels) TG-MS • Allows real time analysis of weight loss and temporal MS identification of evolved gases • Has restricted identification power of the gaseous species emitted according to mass • Is designed for high sensitivity (pg level) and versatility • Permits industrial problem solving capability • Is a real asset in the analysis world, yet by many people still much ignored High Resolution TG-MS • Designed for improved resolution and identification • Allows for stepwise and quasi-isothermal modes TG-MS/MS • Permits complex compositional analysis of thermally activated processes • Allows rapid separation and identification of co-evolving compounds • Differentiates between co-evolving isobaric ions • Stands for real time analysis

FTIR methods which use heatable gas cells, e.g., fast thermolysis FTIR [111], the on-thespot TG-FTIR technique, which avoids transfer lines, monitors the gaseous atmosphere directly above the sample pan; spectral information is obtained which is directly correlated to the recorded mass change as a function of time and temperature. At the end of a TG

8

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… experiment an Evolved Gas Profile (EGP) can be reconstructed from the stored interferograms according to Gram-Schmidt [112]. Each point in this EGP corresponds to an IR spectrum of the evolved components in the TG equipment. Specific Gas Profiles (SGP) and Functional Group Profiles (FGP) can also be reconstructed from the stored interferograms in selected wavenumber windows to detect components with specific group frequencies. Comparison of EGP as a function of time with the DTG curve yields a direct comparison of TGA and spectroscopic data. The detection limits are in the sub μg/s range and dependent on the extinction coefficient of the evolved components [29]. The commercial availability of FTIR systems capable of highly sensitive detection and completely automated sampling and data manipulation, have given TG-FTIR polymer analysis all the advantages of classic IR spectroscopic interpretation. Advantages of FTIR detection are: functional group identification and specific compound analysis on the basis of vibrational spectra; simultaneous spectral information on many species; continuous scanning of effluent from direct thermal processing; quantitative analysis (± 10%) using proper calibration from well-known absorption coefficient information for IR-absorbing compounds [113]; reference (vapour-phase) spectral libraries [114]. In general, the detection limits for components in the condensed phase are a factor-of-ten lower than those in the gas phase. The technique is especially useful for smaller molecules where the high specificity of strong IR absorption bands makes up for the relatively low sensitivity of IR detection. It is rather difficult to use IR to analyse mixtures of compounds with similar functional groups or mixtures of weak IR absorbers in the presence of strong absorbers. The limitation of FTIR lies in detecting only non-symmetrical gas molecules. FTIR cannot detect gases without IR absorbance, e.g., O2, N2, and does not readily distinguish hydrocarbons above C3H8. If the calculated weight loss of observed gases is lower than that measured by TGA, then it can be inferred that other gases are being evolved that are FTIR blind. FTIR uses much lower excitation energy than MS and can therefore detect larger functional groups in evolved gases from TG experiments, such as high-boiling oligomers and heavy tar products which can be analysed as fine aerosols in a gas stream [20, 100]. TG-FTIR also distinguishes structural isomers [115]. Spectral subtraction and spectral search aid the identification of evolved gases, which are often a mixture of products. Nevertheless, for unambiguous identification of unknown volatiles more powerful methods are required. Jansen and co-workers [86] have incorporated a parallel mass spectrometer onto the FTIR stage of a thermogravimetryFTIR (TG-FTIR). The sample is thermally decomposed by TGA and the products collected in a Tenax (absorbent charcoal) trap. After desorption, the products are separated by a GC and the sample split, with 99% going to the IR spectrometer and 1% to the mass spectrometer.

9

Spectroscopy of Rubbers and Rubbery Materials TG-FTIR has become a quite popular, versatile, cost-effective and informative instrument for modern polymer analysts concerned with compositional analysis and degradation/ reaction mechanism studies. The growth rate of TG/FTIR instrumentation currently exceeds that of TG-MS. Simultaneous TG-MS is a very powerful hyphenated technique combining the direct measurement of weight loss as a function of temperature with the use of a sensitive spectrometric detector. The usefulness of MS to TG coupling was suggested already in 1965 [8, 10]. The coupling with MS (in particular time-of-flight MS (ToFMS) and quadrupole mass spectrometry (QMS)) adds chemical analytical features allowing the chemist to assign the evolved gases to the detected weight losses thereby correlating chemical information with the thermal event. Mass spectrometry permits temporal resolution of the gases that are evolved during the thermal or thermo-oxidative degradation of a polymer in controlled atmospheric conditions. It is obvious from the history of TG-MS [25] that the interface is of crucial importance and poses several problems. Within the interface, conditions are converted from the high temperature and (usually) atmospheric pressure of TG to the room temperature and (usually) high vacuum conditions in the mass analyser. Kaisersberger and co-workers [96] have given an excellent account of the basic requirements and practical aspects of coupling of gas analytical methods with thermal analysis instruments. Bart and co-workers [25] have critically discussed a broad variety of TG-MS instrumental design solutions, which depend partly on the sample characteristics and the desired conditions of the thermal degradation, particularly in relation to polymer characterisation. Advantages and disadvantages of TG-MS with respect to other evolved gas analysis techniques are outlined. Mass spectrometry coupling can be achieved by connecting a capillary at the end of the gas-flow system of a thermobalance or by means of a skimmer coupling, integrated into the furnace of the thermal analyser [96]. Although capillary systems have been described as a great limitation of coupling thermal analysis instruments with the mass spectrometer [49] because of presumed restrictions in heating of capillary couplings to a maximum of 200 °C and of most transfer lines to about 250 °C, a transfer experiment with the commercial UV-absorber Tinuvin 234 has shown sufficient sensitivity and transient response of a capillary TG-MS interface for molecular masses up to m/z 450 [116]. Although most TG-MS couplings have concentrated on evolved gas analysis, a few reports concern residue analysis. Wiedemann and co-workers [117] have described a thermomolecular beam analysis (TMBA) technique, which allows to continuously and simultaneously follow the weight of a sample (TG), the composition of its gaseous reaction products (MS) and solid reaction products {high temperature X-ray diffraction (XRD)}. Raemaekers and co-workers [26] have advocated video-imaging (VI-TG-MS). As many of the curves from TGA or DSC are in fact much more complex than might appear at first sight, TG-DSC/DTA with evolved gas analysers (MS and FTIR) are useful 10

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… tools [118]. In a top-loading STA analyser combined with a quadrupole mass spectrometer by means of a two-stage gas inlet system, as described by Kaisersberger and co-workers [50, 63], evolved gases and vapours are collected just above the sample crucible by an orifice. A skimmer is arranged in the compression zone behind this orifice to achieve a parallel molecular beam into the electron impact ion source of the MS. The optimised orifice system and short transfer path in this advanced solution effectively eliminate many of the coupling problems observed in most other designs. The arrangement excels by high sensitivity [49]. A simultaneous TG-DSC-MS capable of operation up to 2000 °C has been reported [119]. As a further extension of the classical TG-MS commercial TG/ DTG/DTA-MS (STA-MS) has been developed (sample size up to 5 g) as well as a macro STA/MS/GC-MS (sample size 500 g), which allows examination of heterogeneous materials and trace analysis [120]. The strengths of MS and FTIR complement each other. MS requires special high vacuum capabilities and more stringent operating conditions but exhibits detection levels which are several orders of magnitude better than FTIR (pg and sub-μg ranges, respectively). In some conditions, MS results can be misleading because of secondary products resulting from ion fragmentation [121]. Yet, MS identifies each individual compound and not a class of compounds of the same functional group characteristics. Both MS and FTIR need the support of a full spectrum vapour phase library. Despite the utility of these techniques, a distinct disadvantage is that very low concentration components may be masked by higher concentration interferants; in such cases, devices based on thermal desorption (absorbent packing, static headspace sampling, cryogenic trapping) may prove useful. More complex thermal decomposition processes may require identification instead of verification, for which single-step hyphenated techniques are ill equipped. Multihyphenated methods have been described which combine a number of functions, such as separation and identification (in various modes). If a pattern of complex species is evolved during heating (as is easily the case for polymers), it is advantageous to achieve separation prior to entering the final phase of the mass spectrometer. This can be accomplished through the use of trapping TG-GC-MS [86, 122 or direct TG-GC-MS [23, 123]. As mentioned already, this analysis is lengthy: for each point on the TG curve (time, temperature) a GC-MS analysis has to be performed. McClennen and co-workers [20] have shown the enhanced capability of the doubly hyphenated thermogravimetrybased analytical techniques TG-IR-MS and TG-GC-MS for detection of minor products. Arii and co-workers [124] have used an integrated simultaneous TG-DTA/GC-MS system in high resolution mode. Controlled rate thermal analysis enhances accuracy of identification and quantification. The TG-DTA/GC-MS system can be operated in two interface modes, namely continuous sampling (direct coupling mode) and intermittent

11

Spectroscopy of Rubbers and Rubbery Materials sampling (trap coupling mode). In the direct coupling mode the GC is bypassed. Intermittent sampling is ‘off-line’. Jansen and co-workers [97] have described on-line TD-GC-FTIR. In the TD-GC-FTIR system a thermal desorption (TD) cold trap injector is used for the temperature-controlled outgassing of the samples with a maximum temperature of 350 °C. The volatile components are transferred to the cold trap by the carrier gas and preconcentrated. After completion of the outgassing process the cold trap is heated very quickly, causing on-column injection of the trapped components onto the gas chromatograph. The technique has recently been extended to include an ion-trap MS. On-line thermal desorption-GC-FTIR-MS can be operated both in parallel and tandem FTIR-MS configuration, where parallel FTIR-MS operation is preferred for several reasons. Compared to FTIR alone, the parallel configuration of FTIR and MS enhances and facilitates the elucidation of the evolved species and furthermore lowers the detection limits from ppm to (sub) ppb level [86]. With a thermal desorption unit capable of accommodating Tenax absorption cartridges sampling with preconcentration of thermogravimetric off-gases in possible which may be followed off-line by analysis of the trapped components via TD-GC-FTIR-MS [86]. A major disadvantage in routine analysis by these techniques is the throughput of samples; the analysis time is determined by the longest retained compound. Quite obviously, neither TG-GC-MS [20, 22, 23, 87, 125-129] nor TG-GC-FTIR [20, 24] provide continuous monitoring of evolved gases. Other significant disadvantages of multihyphenated systems are complexity, cost, and the need for a trained operator. It is therefore not at all certain that such techniques hold the future.

1.2.1 Experimental

1.2.1.1 Sample Requirements for Polymers For TG work it is generally recommended to use as little sample as possible within the limits of resolution of the microbalance. The homogeneity of a sample can sometimes limit how little sample can be used, e.g., in case of polymeric materials. Powdered samples, of small particulate size, have the ideal form for TG studies. However, in polymer science, samples are often also in the form of films, fibres, sheets, pellets, granules or blocks. The packing density should be as uniform as possible. Multicomponent separation of a rubber material performed with TG is typically as follows: rapid heating in inert atmosphere (nitrogen) up to 100 °C (for loss of volatile oils and 12

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… extenders), successive heating up to 600 °C (for decomposition of the rubber component), heating in oxygen to 950 °C (for combustion of carbon black) and determination of the residue (fillers). For (multi) hyphenated thermogravimetric analysis techniques typically a few mg of sample are sufficient. Sample requirements for TG-MS vary according to the amount of volatile substances liberated and the sensitivity of the mass spectrometer for a given component. Moreover, as the temperature that corresponds to the maximum evolution of a fragment is particle-size dependent, there is a size requirement as well. When determining low level additives it is often quite necessary to weigh out as much as 100 mg; sample sizes of 3-6 mg are a usual acceptable compromise between efficient balance operation and overloading the mass spectrometer. For sample sizes < 5 mg, problems associated with heat transfer and surface collisions are minimised; these samples are more amenable to maintaining a low system pressure and reducing bimolecular collisions. The pressure limit condition for the mass spectrometer implies that there is an upper limit for the product of the rate of heating, dT/dt, and the mass M of the compound to be decomposed: dT/dt. M < C, where the value of the constant C is dependent upon instrument and sample. This condition should especially be considered in case of quantitative measurements [130].

1.2.1.2 Calibration and Interlaboratory Reproducibility Accurate temperature calibration using the ASTM temperature standards [131, 132] is common practice for DSC and DTA. Calibration of thermobalances is more cumbersome. The key to proper use of TGA is to recognise that the decomposition temperatures measured are procedural and dependent on both sample and instrument related parameters [30]. Considerable experimental control must be exercised at all stages of the technique to ensure adequate reproducibility on a comparative basis. For (intralaboratory) standardisation purposes it is absolutely required to respect and report a number of measurement variables. ICTA recommendations should be followed [133-135] and should accompany the TG record. During the course of experiments the optimum conditions should be standardised and maintained within a given series of samples. Affolter and coworkers [136] have described interlaboratory tests on thermal analysis of polymers. Courtault [137] has critically considered calibration of TG, MS and the TG-MS coupling. The TG (-MS) apparatus is often calibrated with calcium oxalate monohydrate. Bart and co-workers [25] have discussed critically the calibration and interlaboratory reproducibility of TG-MS. The major hurdle to be overcome before a technique gains acceptance within the analytical community and in practice is the achievement of interlaboratory validation. There still 13

Spectroscopy of Rubbers and Rubbery Materials appears to be sufficient room for improvement in interlaboratory reproducibility for most other multihyphenated techniques considered in this paper.

1.3 Applications of Hyphenated Thermogravimetric Analysis Techniques Thermogravimetric analysis can yield a considerable amount of information about the composition of an elastomer system and is a common means of testing of EPDM rubber compounds [51]. The DTG curve may serve as an identifier of the type of elastomer in a compounded formulation. Sircar [138] has reviewed the analysis of elastomer vulcanisate composition by TG/DTG techniques. The classical ASTM method, D297-93 [139], is too lengthy to be of much practical use on a routine basis, often requires preliminary identification of the polymer and is costly. TG has gained itself wide acceptance as a method for quantitative compositional analysis of vulcanisates: ASTM E1131 [140], is basically designed for the analysis of rubber compounds [141]. Thermogravimetric analysis can be used to determine: •

water, solvents, monomers and other volatile substances

•

oil, extenders, plasticisers and other high temperature volatile substances

•

elastomer, organic fillers by pyrolysis

•

carbon black and graphite by high temperature oxidation

•

carbonates, e.g., CaCO3, by decarbonation at elevated temperatures

•

ash, pigments and inert fillers as measured residue

•

thermal stability, kinetics of decomposition and thermal endurance.

TGA can be used as a quality control method in the production of rubber compounds (ASTM Task Group E 3701.09), and is a recognised troubleshooting tool in the rubber industry. Classical analysis of rubber compounds for identification of ingredients such as accelerators, antioxidants, and antiozonants requires a very difficult separation step of such additives from the rubber compound prior to analysis. Vulcanised rubbers are even more difficult to analyse than thermoplastic polymers due to the compounding ingredients being locked into the matrix by carbon black and cross-linking of the rubber polymer. The classical method for separating volatile additives from rubber is by solvent extraction. TGA measurements can

14

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… differentiate between the most volatile components in rubber compounds, for example moisture, plasticiser, residual solvent, processing oils and extenders, or other low boiling point (below 300 °C) materials and components of medium volatility including the elastomer portion of the compound. Although TG/DTG is a potentially effective analytical technique for the compositional analysis of compounded elastomers such analyses are not always completely accurate, i.e., a 100% materials balance cannot always be obtained [28]. This may be due to an overlap of low molecular weight volatile material, e.g., oils, plasticisers and resins, with polymer decomposition products, the decomposition of polymer blend components in a similar temperature region, the formation of char, which decomposes in the region assigned to carbon black, or the carry-over of early stage decomposition products to the ash region. Typically, sulphur, accelerators, antioxidants and antidegradants are not observed as independent weight loss in TG curves. For many systems weight losses overlap for most choices of method parameters. Consequently, high resolution or reduced pressure methods are frequently used. Apart from incomplete analysis due to poor resolution and the inability to separate oligomers of polymer and oil, desorption methods such as TGA may produce volatile fragments of additives. The interferences are discussed by Sircar [138]. Thus, there are practical limits to the kind and degree of information that can be extracted from TG/ DTG analyses of unknown vulcanisate compositions. Table 1.4 shows that the most numerous applications of hyphenated thermogravimetric techniques for the study of elastomeric material make use of TG-DTA, followed by TGFTIR, TG-DSC, TG-MS and TG-DTA-MS, with only occasional recourse to TG-DSCMS and TG-GC-MS. Table 1.5 indicates the general performance characteristics of the thermogravimetric techniques in use for the study of elastomeric materials. There is a clearly defined industrial need for reliable and fast methods to study the stabilisation and characterisation of elastomeric materials, to obtain knowledge on their properties and behaviour, the effects of modifying structure, additives and processes to engineer the most suitable material for any given application. Also, there is an increasing (environmental) need to study the effects of ageing, the thermal stability, the degradation processes and the products of decomposition under a wide range of conditions. Evolved gas analysis may be used in two modes: •

as a stand-alone technique, such as temperature-programmed reduction

•

as a hyphenated technique (TG, DSC, or STA combined with FTIR or MS).

Hyphenated thermogravimetric techniques have been used in a wide variety of qualitative and quantitative problem-solving cases for elastomeric materials:

15


Table 1.4 Hyphenated thermogravimetric analysis studies of elastomeric materials TG-FTIR

Vulcanisation [32], ageing characterisation [39, 48], sulphur components in rubber [31], polyurethanes [37], polymer degradation mechanisms [30, 40, 41], identification of base polymers [36, 43, 44], thermal stability [46], grafted flame retardants [47], differentiation of EVA rubbers [45] and AN-NBR rubbers [36, 44], degradation of chlorinated natural rubber [42].

TG-MS

Kinetics of degradation [54], ageing characterisation [48], quality control [51], controlled release [52], additive analysis [88, 50, 142, 143, 144], outgassing [48, 49, 50, 145, 146], thermal degradation [25, 53, 54, 147], product development [148].

TG-DSC

Additive interactions [58], reactivity [59-62].

TG-DSC-MS

EPDM and chlorinated rubbers [63].

TG-DTA

Characterisation of carbon black [149], flammability evaluation [64], polymer degradation studies [65], ageing studies [70-72], product control [77, 81], combustion performance [83], safety evaluation [83], antioxidation activity [68], pyrolysis of rubbers [82], thermal stability [67, 69, 76, 77], interfacial junctions in viscoelastic composites [78], weathering [72], vulcanisation [73], oxidative behaviour [79], materials evaluation [80], failure analyses [81].

TG/DTG-DTA-FTIR

Weathering [39].

TG-DTA-MS

Combustion performance [59-62, 85], safety evaluation [59-62, 85].

TG-GC-MS

Co-processing of waste polymers [150].

TG-Tenax-TD-GC-FTIR-MS

Outgassing [86, 175].

AN: acrylonitrile

16


Table 1.5 Comparison of hyphenated thermogravimetric techniques for elastomer analysis Methoda

Thermal effects

Evolved gas analysis

with mass change

withoutb mass change

quantitative

qualitativee

TG

+

-

-

-

TG-FTIR

+

-

+c

+

TG-MS

+

-

+c

+

TG-DSC

+

+

-

-

TG-DSC-MS

+

+

+c

+

TG-DTA

+

+

-

-

TG-DTA-MS

+

+

+c

+

TG-GC-MS

+

-

+d

+

a

All combined techniques measure various effects on the same sample Energetics of phase transitions or reactions (melting, crystallisation, TG, cross-linking, etc.) c With appropriate calibration the system is suitable for quantitative determination. d Can be made quantitative e Structural assignment on the basis of molecular ion mass, fragmentation pattern or vibrational mode b

a. Thermal stability and degradation studies • • • • •

ageing characterisation in thermal and thermo-oxidative stability studies weathering identification of degradation products elucidation of polymer degradation mechanisms intercomparison of materials properties characterisation of insoluble polymeric materials.

b. Structural characterisation and chemical analysis • • •

identification of base polymers and components differentiation of materials fingerprint identification (classification of competitor products).

17

Spectroscopy of Rubbers and Rubbery Materials c. Reactivity and curing • •

vulcanisation characterisation of cure processes.

d. Product development • • • • • •

optimisation of production processes of elastomeric materials product quality and batch-to-batch control determination of end-use conditions performance behaviour failure analysis product recycling.

e. Analysis of evolved gases during synthesis, processing and recycling • • • • • • • • •

outgassing phenomena trace solvent, monomer inclusions or impurity analysis analysis of additives or processing agents additive interactions determination of the effect of stabilisers controlled release combustion performance health protection studies toxicology of combustion.

f.

Kinetics

• •

kinetics of degradation life-time predictions (oxidative induction time; OIT).

g. Quantitative analysis •

determination of the composition of elastomers.

Quantification of sample components by conventional TG may be made only if a stoichiometric relationship exists between the gaseous decomposition products and the original material. TG-FTIR allows quantitative analysis to be performed (calibration necessary) even when more than one component of interest pyrolyses during a single weight loss [152]. Courtault [137] has described quantitative aspects of TG-MS coupling, which is still a difficult technique. Dyszel [153] has determined the styrene content in styrene-butadiene

18

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… rubbers. The complexity of quantitative measurements of thermal analysis using mass spectrometry has recently been discussed by Morelli [154]. The development of a Hi-Res TG-MS [17] is beneficial to the ability to obtain more accurate, quantitative results. Quantitative interpretation of TG-MS data requires calibration of the system, i.e., the determination of the relationship between the observed intensities of the ion currents and the amount of the analysed species. Quantitative aspects of MS couplings have recently been treated very clearly by Maciejewski and co-workers [155, 156], also introducing a new experimental technology, Pulse Thermal Analysis (PTA). PTA enables the introduction of a well-defined amount of a gas to the system at any temperature (non-isothermal) and/or time (isothermal mode). Injected pulses can be used as a reference for the quantification of the signals originating from the evolution of gas(es) formed during the decomposition of solids. A linear dependence between the amount of injected gas and the intensity of mass spectrometry signals enables the quantification of mass spectroscopic data. An accuracy for the evolved species below 0.01 wt.% has been quoted. Unless results are made quantitative it is arduous to draw any definite conclusions about mechanisms of decomposition. Table 1.4 lists applications of various hyphenated thermogravimetric techniques for the study of elastomeric materials.

1.3.1 Thermogravimetry - Fourier Transform Infrared Spectroscopy The applications of simultaneous TG-FTIR to elastomeric materials have been reviewed in the past. Manley [32] has described thermal methods of analysis of rubbers and plastics, including TGA, DTA, DSC, TMA, Thermal volatilisation analysis (TVA), TG-FTIR and TG-MS and has indicated vulcanisation as an important application. Carangelo and coworkers [31] have reviewed the applications of the combination of TG and evolved gas analysis by FTIR. The authors report TG-FTIR analysis of evolved products (CO2, NH3, CH3COOH and olefins) from a polyethylene with rubber additive. The TG-FTIR system performs quantitative measurements, and preserves and monitors very high molecular weight condensibles. The technique has proven useful for many applications (Table 1.6). Mittleman and co-workers [30] have addressed the role of TG-FTIR in the determination of polymer degradation pathways. Some general applications of TG-FTIR are evolved gas analysis, identification of polymeric materials, additive analysis, determination of residual solvents, degradation of polymers, sulphur components from oil shale and rubber, contaminants in catalysts, hydrocarbons in source rock, nitrogen species from waste oil, aldehydes in wood and lignins, nicotine in tobacco and related products, moisture in pharmaceuticals, characterisation of minerals and coal, determination of kinetic parameters and solid fuel analysis.

19


Table 1.6 Polymer chemistry applications for TG-FTIR after [33] Solvent and water retention Curing reactions Vulcanisation reactions Isothermal ageing Product stability Thermal degradation Identification of processing aids Plasticisers Mould lubricants Blowing agents Antioxidants Flame retardants Safety concerns Product safety Product liability Fire hazards

The potential applications of an integrated TG-FTIR system were discussed by various authors [33, 34]. Schönherr [35, 36] has discussed the advantages of the method, in particular for rubber analysis. TG-FTIR examination of a polybutadiene sample with a high proportion of inorganic fillers and spectral subtraction procedures identified water and plasticiser at 200 °C; CO2, CO, H2O, methane, ethylene, n-butane, n-pentane and cyclic hydrocarbons at 500 °C. The results indicate that a single sample weight loss may well correspond to a very complex mixture of evolved gases [57]. Another synthetic rubber sample evolved CO, CO2, CS2 and hydrocarbon traces (at about 250 °C), methane and short-chained alkenes (between 375 °C and 540 °C) [33]. TG-FTIR of a gold-filled epoxy resulted in evolution of 1-epoxy-7-oxabicyclo [4.1.0] heptane at 185 °C, carboxylic acid and CO2 at 330 °C, CO2, phenol and carboxylic acid between 360 °C and 460 °C [33]. TG-FTIR of a silicone O-ring under nitrogen showed formaldehyde and hexamethylcyclotrisiloxane at 450 °C, and a hexamethylcyclotrisiloxane octamethylcyclotetrasiloxane mixture up to 550 °C [33]. Brame [37] used TG-FTIR for 20

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… the determination of minor components in polyurethanes, which could not be discerned by FTIR without a thermal separation procedure. The technique was also used for the study of the decomposition of polybutadiene [38] and of weathered sealants on a silicone and polyurethane basis (in an STA-FTIR configuration) [39], as well as for the study of the thermal degradation mechanism of tetrafluoroethylene-propylene copolymers [40]. Although the exact mechanism for the thermal degradation of Aflas FA-150C rubber could not be defined, comparative TGFTIR experiments (see Figure 1.3) demonstrated that the main degradation products (see Figure 1.4) were mixed hydrocarbon/fluorocarbons similar to those observed in Tefzel 200 and 280 (alternating copolymer) degradation rather than what would be expected if degradation occurred by unzipping to monomers. Suzuki and co-workers [41] have studied the thermal degradation of acrylonitrile-butadiene-styrene (ABS) terpolymer by TG-FTIR (see Figure 1.5) and compared with that of polystyrene (PS), butadiene rubber (BR), styrene-acrylonitrile copolymer (SAN) and polyacrylonitrile. TGFTIR has also been used for a study of the thermal-oxidative and thermal degradation processes of chlorinated natural rubber (CNR) [42]. Thermal degradation of CNR is a one-step reaction with dehydrochlorination taking place at 160-190 °C, whereas thermooxidative degradation of CNR is a two-step process with production of HCl and CO2 at 160-390 °C and of CO2 at 390-575 °C.

Figure 1.3 Evolved gas profiles (EPG) for PTFE, Aflas and PP upon heating at 10 °C / min. Ordinate axis values are arbitrary and simply relative Reproduced with permission from H.G. Schild, Journal of Polymer Science A, 1993, 31, 6, 1629. Copyright 1993, John Wiley and Sons

21


Figure 1.4 Comparison of Aflas degradation product spectra at moderate (531 °C) and high (726 °C) temperatures Reproduced with permission from H.G. Schild, Journal of Polymer Science A, 1993, 31, 6, 1629. Copyright 1993, John Wiley and Sons

Figure 1.5 Infra red spectrum of the gases evolved from ABS terpolymer at 400 °C Reproduced with permission from M. Suzuki and C.A. Wilkie, Polymer Degradation and Stability, 1995, 47, 217. Copyright 1995, Elsevier Science Ltd

22

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… Schönherr [43] has described the combination of decomposition in a thermogravimetry oven and FTIR spectroscopy for the identification of base polymers in elastomers, as exemplified for nitrile rubber, and has presented infrared spectra for decomposition products of various rubbers. The same author [36] studied use of the integrated TGFTIR system for the identification of sixteen vulcanised rubbers in mechanical goods reporting the characteristic infrared spectra of the degradation products at temperatures ranging from 334 °C to 635 °C. As is well known, the properties of technical rubber goods are determined to a great extent by the polymer type and comonomer ratios. Identification of the base polymer is therefore of considerable interest, in particular in damage cases, raw material and product quality control. Schönherr [44] used TG-FTIR to analyse an acrylonitrile(AN)-butadiene rubber (NBR). TG-FTIR of NBR shows CS2 at 270 °C, HCN, NH3 and hydrocarbons at 450 °C. The method was also applied to distinguish various uncured NBR (Perbunan NT and Krynac) types with different AN content (18 to 48%). EGP and three FGP (characterising HCN at 700-730 cm-1, NH3 at 800-1160 cm-1 and hydrocarbons at 2800-3000 cm-1) were reported. NBR rubbers with different AN content are best distinguished on the basis of the absorption intensities of HCN (714 cm-1) and CH compounds (2900-3000 cm-1). It thus appears that the technically most important elastomer types can be discriminated by IR spectra of their gaseous decomposition products [36, 44, 157]. The integrated TG-FTIR system has been similarly used to analyse uncured and vulcanised ethylene vinyl acetate copolymers (EVA) with different vinyl acetate (VAC) contents (40-70%) [45]. Polyesterurethanes, polycarbonate and silicone rubbers have been studied by TG-TenaxFTIR/MS. The degradation of polyesterurethanes yields CO2, water, tetrahydrofurans, cyclopentanone, dicarbonic acid, and aliphatic diols and esters. The thermal decomposition of silicone rubbers leads to the formation of polychlorinated biphenyls which are produced in small amounts and can be observed in the mass spectrometer [86]. Post and co-workers [49] have used TG-FTIR to study the outgassing of a plasticiser (type and amount) from an ethylene-propylene-diene terpolymer (EPDM) compound. Figure 1.6 shows the thermogravimetric decomposition behaviour of the EPDM compound. The plasticiser emerges in the first mass-loss step, which was identified as adipic acid diisobutylester by on-line infrared analysis. Melissaris and co-workers [46] have studied various novel p-ethynyl-substituted rigid rod monomers by DSC, TG and TG-FTIR and identified the main decomposition products. Void-free neat resin mouldings were made by compression moulding the monomers followed by heating. Wilkie and co-workers [47, 158] have used TG and TG-FTIR to study grafting of vinyl monomers, such as methacrylic acid, onto styrene-butadiene block rubber (SBS) or ABS

23


Figure 1.6 Thermogravimetric decomposition behaviour of an EPDM compound Reproduced with kind permission from E. Post, S. Rahner and F. Giblin, Proceedings of Antec ‘97, Toronto, Canada, Volume 2, 2300, Figure 7. Copyright 1997, Society of Plastic Engineers [49] to achieve flame retardancy. Little difference was observed in the degradation pathway between the grafted and virgin polymers. On-the-spot TG-FTIR of poly(butylene terephthalate)(PBT)/octabromodiphenylether (molecular weight (MW) 801 Da) revealed the brominated diphenylether at 275 °C and terephthalic acid (the starting monomer of PBT) at 425 °C [110]. Similar high molecular species have never been reported in TGMS experiments; the flame retardant was also missed out in off-line TG-GC-FTIR-MS analysis. The same authors [29] used the technique for the study of a carbon black-filled styrene-butadiene rubber (SBR) providing information about the composition, including organic additives, polymers, carbon black and inorganic fillers (see Figure 1.7). At 250 °C, water, CS2 and morpholine were detected. The latter two components are degradation products of the accelerator used, 2-(morpholinothio) benzothiazole. In an examination of an ABS/polycarbonate (PC) blend with 8% triphenylphosphate (TPP), in addition to the EGP, the SGP for the specific wavenumber windows of TPP (900-1200 cm-1), aromatic compounds (3000-3100 cm-1), and carbon oxides originating from PC (2200-2300 cm-1) were obtained. TPP evolving first was detected at about 150 °C (detection limit 0.5 μg/ s) [29]. With the direct IR detection of the on-the-spot technique, there is no loss of evolved components by cold spots or discrimination of high-molecular-mass components. In addition to the chemical composition of the evolved components, the technique also provides information on the sequence and kinetics of the mass-loss process, which may not show up in the mass-loss curve. TG-IR has also been used to examine the thermally induced decomposition products of polyvinyl chloride (PVC), polyacrylamide, tetrafluoroethylene-propylene, styrene-

24


Figure 1.7 IR spectrum of the degradation products of the accelerator of carbon blackfilled SBR at about 250 °C Reproduced with kind permission from J.A.J. Jansen, J.H. v.d. Maas, A. Posthuma De Boer, Applied Spectroscopy, 1992, 46, 1, 88. Copyright 1992, the Society for Applied Spectroscopy butadiene composites and ethylene vinyl acetate (E/VAc) copolymers [159]. Because of the mixture of products generated, identification of the constituents could only be made when small molecules with simple infrared spectra, such as CO, CO2 and HCl were generated. Alternatively, the effluent of TG-IR is captured on a trap constructed from a GC capillary injector liner with Tenax solid phase adsorbent and analysed by GC-IR. Other reported TG-FTIR applications concern: polyesterurethane [160], silicone rubber [86, 33], ABS [161], synthetic rubbers used in the automotive industry [38] and the determination of butadiene and styrene in unknown composites [162].

1.3.2 Thermogravimetry-Mass Spectrometry Bart and co-workers [25] and others [34, 101, 163] have reviewed the application of TGMS for the study of polymeric materials, thermoplastics, thermosets and elastomers. This thermoanalytical technique is used for the structural characterisation of homopolymers, copolymers, polymeric blends and composites and finds application in the detection of monomeric residuals, solvents, additives, (toxic) degradation products, etc. Information is

25

Spectroscopy of Rubbers and Rubbery Materials also obtained on the mechanism of solid-state reactions, chemical reactivity and curing; TG-MS is also beneficial in matters of product formulation and development. TG-MS is an ideal technique for identifying residual volatiles in polymers. The detection of residual volatiles (and of other impurities) can often yield clues as to manufacturing processes. In many cases, such as in the determination of highly volatile materials, of residual solvents or plasticisers, use of TG-MS is requested. Specifically, there are reports on the entrapment of curing volatiles in bismaleimide laminates [145] and elastomers [48], on the detection of a curing agent (dicumylperoxide) in EPDM rubbers and of bromine flame retardants in electronic waste [50], of plasticisers such as bambuterol hydrochloride [142] or TPP and diethylterephthalate in cellulose acetate [143], on solvent extraction and formaldehyde loss in phenolic resins [164], and on the evolution of toxic compounds from PVC and polyurethane foams [146]. Kaisersberger and co-workers [50] have reported detection of nitrosamine precursory compounds during rubber vulcanisation (originating from vulcanisation agents) and the determination of toxic or environmentally damaging exhaust gases during technical burning processes (polycyclic aromatic compounds, PCB, etc). TG-MS couplings are increasingly used by the rubber industry especially since aromatic plasticisers are toxicologically suspect. Post and co-workers [49] used the skimmer-MS coupling in TGMS measurements to study the outgassing of a plasticiser from an EPDM compound. The sensitivity needed for identification of additives is, in general, greater than that needed for identification of volatile pyrolysates. In other applications the pattern of evolution of styrene, butadiene and acrylonitrile as a function of temperature has provided a unique way for classifying different types of ABS. The loss of the antioxidant butylated hydroxytoluene (BHT) was also detected by MS preceding EVA copolymer degradation [165]; BHT was identified at a concentration level of 20 ppm. Lehrle and co-workers [52] have described a successful controlled release system for the stabilisation of rubber by encapsulating efficient but rather mobile antioxidants to prevent loss from the host polymer. The performance of the controlledrelease of the antioxidant BHT from alginate matrix particles was studied by means of DSC, TG and TG-MS. Polyisoprene rubber is more resistant to oxidation when protected in this way than by the equivalent concentration of unencapsulated antioxidant. TG-MS is ideally suited to reveal differences in pressure behaviour during thermal decomposition of materials. This has been illustrated by Mol [144] in TG-MS analysis of toluene diisocyanate (TDI) and methylene bis-4-phenol isocyanate (MDI)-based polyurethanes, where the observed greater increase in pressure for the TDI polyurethane than for the MDI derivative indicates a higher loss of low molecular weight fragments. This is not possible to deduct from the TG curves alone. Such indications are of great

26

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… importance when evaluating kinetics or considering mechanically similar materials for heat sensitive applications. The presence of impurities and additives such as PVC as a fire retardant in polyurethanes has also been detected by means of TG-MS analysis [144]. PU foam containing the flame retardant tetrakis (2-chloroethyl) ethylenediphosphate decomposes in an oxidative atmosphere at standard pressure in one rapid reaction whereby several highly toxic species are formed; the TG-MS detection limit of this flame retardant was determined [166]. The most striking fact is the very early generation of vinylchloride from the flame retardant, while HCN is produced at much higher temperatures. Bart and co-workers [25] have reported the degradation of a complex, urethane-based, thermoplastic elastomer and have identified the volatile products (mainly CO2, 1,3butadiene, tetrahydrofuran (THF) and cyclopentanone) by TG-MS and GC-MS. While CO2, THF and HCN were observed by TG-MS a much richer evolved gas composition with m/z = 44, 54, 70, 72, 84, 90, 200 (all components identified) was detected with TGCT-GC-MS. The main degradation mechanisms were identified as thermal degradation of ester bonds of adipic acid and urethane moieties. This very complex polymer material was chosen deliberately in an attempt to illustrate the limitations of the TG-MS technique. As expected, not all the gases were readily identified. Several of the weight loss transitions are caused by complex decomposition processes resulting in multiple materials being evolved. Particularly if some of the components are present in minor amounts, the separation becomes an integral part of the analysis. Recourse to GC-MS is an obvious choice in case of complex mixtures in the TG atmosphere for which simple MS identification of the component organic substances is not possible [88, 167]. Also several EPDM products were studied by means of TG-MS [25]. TG/DTG of an EPDM without filler and plasticiser shows that during the maximum weight loss phenomenon ENB (m/ z = 66, 91, 105 Da), aliphatics (m/z = 43, 56, 69 Da) and olefines (ethene: m/z = 26, 27 Da; propene: m/z = 40, 41, 42 Da) are detected. The dynamic DTG and MS curves in inert atmosphere of an EPDM compound charged with oil, filler and carbon black, indicates loss of oil (maximum at 336 °C), thermal stability of the polymer up to about 420 °C (maximum decomposition at 485 °C), and decarboxylation of the filler at 730 °C (CO2: m/z = 12, 44 Da); finally, above 900 °C in O2 atmosphere carbon black is detected. The same authors [25] have reported a TG-MS study of EPDM-SBR blends. Griffiths and co-workers [53] have described the effect of the binder Alloprene (a chlorinated rubber) on the ignition characteristics of some pyrotechnic compositions. Sklemin and coworkers [54] have studied the kinetics of thermal degradation and evolution of volatile products during thermolysis of ethylene-propylene rubber, synthetic fibre and carbon-fibre reinforced phenol-formaldehyde copolymers. The temperature of initial degradation (5 wt.% loss) shifted to higher temperatures with increasing heating rate. Degradation rate,

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Spectroscopy of Rubbers and Rubbery Materials rate constants of degradation and activation energies of degradation were determined. Other applications of TG-MS in kinetic studies of polymers concern the (oxidative) thermal degradation of polymethyl methacrylate (PMMA), polyurethanes and acrylonitrile (ACN)/ styrene copolymer [147]. Möhler and co-workers [51] have reviewed the suitability and possibilities for thermal analysis in accordance with DIN 51005 [168] for quality control (batch-to-batch control and optimisation) in the rubber industry, with special reference to heat flux DSC, TG, TG-MS and thermomechanical methods. Results are reported for various rubber compounds and blends. TG-ToFMS has been used in product development by Kleineberg and co-workers [148] as a means for the evaluation of the toxicity potential in normal use and catastrophic situations of some 300 fire retardant materials applied in interiors of passenger and cargo aircraft. ToFMS with its inherent high speed scanning capability enables the toxicologist to relate conventional TGA information to the unequivocal identification of potentially toxic thermal decomposition products. TG-ToFMS of a carboxynitroso rubber showed abrupt, complete decomposition at 292 °C. As shown in Table 1.7 the mass data indicated two primary decomposition products, namely carbonyl fluoride (m/z 66, 47, 50, 31, 19), perfluoro-N-methyl methylenimine) (m/z 133, 69, 114, 31, 50, 45, 26, 57, 64, 12, 19) and the secondary reaction products CO 2 (m/z 44) and trifluoromethylisocyanate (m/z 111) and corrosion products (HF, SiF4, (CF3)2NH; m/z 85, etc.). Quantitative determination was achieved through correlation of the mass spectrometric and the thermogravimetric data. TG-MS and TG-FTIR have been used for ageing characterisation of the methylphenyl silicone elastomers GE 566 (containing ferric oxide and silica filler) and GE 567 with the

Table 1.7 Decomposition products from a carboxynitroso rubber m/z values Carbonylfluoride

COF2

66, 47, 50, 31, 19

Perfluoro-N-methylmethylenimine

CF2=N-CF3

133, 69, 114, 31, 50, 45, 26, 57, 64, 12, 19

Trifluoromethylisocyanate

CF3-N=C=O

111

Carbon dioxide

CO2

44

Hydrofluoric acid

HF

Silicon tetrafluoride

SiF4

Perfluorodimethylamine

(CF3)2NH

28

85

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… primary object of identification and quantitation of volatiles in these materials [48]. Cure catalysts were either 0.1 wt.% dibutyl tin dioleate or 0.5 wt.% dibutyl tin dilaurate. TG-MS identified the trimethyl silyl fragment ion, benzene, styrene and cyclic oligomers of dimethyl siloxane. The major difference between GE 566 and GE 567 found by TGMS was the much greater amount of weight loss and the presence of styrene for the latter, irrespective of the catalyst used. Other reported TG-MS applications concern polybutadiene [153], styrene-butadiene rubbers [153], gums [14], polyisoprenes [52], polyurethanes [144, 146, 147, 166], ABS [144], chlorosulphonated polyethylene elastomer [169, 170] and elastomer blends (NBR/SBR/ BR) [13]. Table 1.5 summarises the use of advanced TG-MS systems in elastomer analysis.

1.3.3 Thermogravimetry - Differential Scanning Calorimetry Thermal analysis is a useful tool in the quality control of many incoming routine materials, which can be tested against a reference standard developed internally by analysing a large number of samples of known performance criteria to ensure that the quality of supplies is maintained. Solid elastomers can be identified by glass transition temperature (Tg) [70]. The rubber industry uses thousands of different raw materials, and this number is steadily increasing. These materials are listed in [172]. Hyphenated TG-DSC provides simultaneous measurement of transition temperatures/ heat flows and weight changes in materials, thereby simplifying interpretation, increasing productivity and assuring identical experimental sampling. TG-DSC can be used to study thermal stability, decomposition behaviour, drying and firing processes, reaction rates, sample composition, kinetics of reaction, transition and reaction temperatures, melting and crystallisation processes as well as reaction mechanisms. Although TG is an excellent technique for the compositional analysis of compounded elastomers, it does not reveal the extent of cure. For that purpose DSC is required. Redfern [57] has reported TG-DSC of an uncured polyimide resin in which a more accurate determination of the quantitative measurement of the heat of cure is made possible by the simultaneous technique. Kodama and co-workers [58] have reported TG-DSC curves for the analysis of the interaction between vulcanisation accelerators (tetramethylthiuram disulphide, dibenzothiazolyl disulphide, diphenylguanidine and N-cyclohexyl-2-benzothiazolylsulphenamide) and fillers (carbon black, white carbon, hard clay and CaCO3). The initial melting point (MP) of the accelerators was largely influenced by the fillers. The higher the surface activity of the filler is, the lower and wider the melting range becomes. Emmott and co-workers [59] have investigated the complex reaction between Sr(NO3)2 and the binder Alloprene (a pyrotechnic system) at about 300 °C by simultaneous TG-

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Spectroscopy of Rubbers and Rubbery Materials DSC and TG-DTA-MS. The same techniques were used to examine the Ti-NaNO3-Alloprene system [60, 61]. The NaNO3 reacts with HCl from Alloprene decomposition and the carbonaceous residue of the Alloprene is oxidised by the O and NOx formed in the first stage. Simultaneous TG-DSC and TG-DTA-MS were also used to investigate the first reaction stage in the related Mg-NaNO3-Alloprene (chlorinated rubber) pyrotechnic system at about 300 °C [62]. The reduction in overall weight losses and the amount of main gaseous products evolved (CO2, NO, H2O) suggested interaction of Mg with the gaseous products of the reaction of HCl and NaNO3; this contrasts with the action of Ti [61]. The TG-DSC technique has recently been reviewed [56]. Redfern [57] has reviewed single sample simultaneous thermal analysis, i.e., TG-DSC and TG-DTA studies of polymers.

1.3.4 Thermogravimetry - Differential Scanning Calorimetry - Mass Spectrometry Whereas Redfern [57] has pointed out the advantages of simultaneous thermal analysis techniques (particularly TG-DSC and TG-DTA) over techniques conducted singly, an even more complete thermal profile is provided when a thermal analyser is coupled to some form of gas analyser (MS or FTIR). Möhler and co-workers [51] have reported TG-DSC-MS of the thermal decomposition of the vulcanisation accelerator tetramethyl thiuram disulphide (TMTD) in rubber; degradation of TMTD starts at about 155 °C, as evidenced by m/z 76 (CS2) and 44 (radical of the secondary dimethylamine). The high sensitivity of the instrumental combination was demonstrated by Kaisersberger and co-workers [63] who published TG-DSC-MS data for EPDM showing cumyloxy radicals (m/z 135, 136) from the dicumylperoxide system (see Figure 1.8). Without the MS data, the mass loss in the range from 240 °C to 400 °C would only have been attributed to the plasticiser content. Hyphenation prevents both a misinterpretation of the results and permits optimisation of the process by adjusting the amount of DCP added to the elastomer prior to vulcanisation. TG-DSC-MS was also used to recognise epoxy resin fragments (m/z 58, 92, 135) in electronic scrap from the automobile industry and bromine flame retardants in electronic waste [50, 63].

1.3.5 Thermogravimetry - Differential Thermal Analysis Since TG and DTA complement each other, it is an obvious move to attempt both investigations simultaneously [173]. TG-DTA measures mass and energy changes as a function of temperature or time. Depending on the atmospheric conditions (vacuum, inert or air conditions) thermal or oxidative stability is measured. Typical TG-DTA

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Figure 1.8 EPDM rubber with peroxide curing agent, detection of cumyloxy fragments Reproduced with permission from J. Janoschek, E. Kaisersberger and E. Post, Polymer Materials Science Engineering, 1993, 69, 458. Copyright 1993, the American Chemical Society [63]

applications are thermal and oxidative stability, determination of relative components, decomposition temperatures and thermal decay reactions, action of heat stabilisers, thermal ageing. The main use of DTA is to detect the initial temperatures of thermal processes and qualitatively characterise them as endothermic or exothermic, reversible or irreversible, etc. Ideally, the area under the DTA peak should be proportional to the heat of the process that gave rise to the peak. TG/DTG-DTA instruments are commercially available. Negri and co-workers [149] have applied TG-DTA to the characterisation of different types of carbon black in natural rubber (NR) vulcanisates. The method allowed the determination of the overall carbon black content, but where combinations of different blacks were present it was not possible to determine the proportion of each type, as instead is possible by means of pyrolysis gas chromatography FID (PyGC-FID) [174]. TG-DTA has also been used to correlate thermogravimetric analysis in airflow and N2 gas flow and some other micro-scale flammability tests, i.e., the oxygen index, hot-plate ignition and drum friction tests, on covers of different flame-resistant and non flameresistant rubber conveyer belts [64]. The minimum temperatures at which rapid weight loss of each sample began to appear were determined and compared with the results from the micro-scale flammability tests.

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Spectroscopy of Rubbers and Rubbery Materials Lee and co-workers [65] have studied the thermal degradation of polyetherurethane elastomers based on polytetramethyleneglycol (PTMG, MW 1000 or 2000), 4,4´diphenylmethanediisocyanate and 4,4´-diaminodiphenylmethane, m-phenylenediamine or p-phenylenediamine by means of TG-DTA. Degradation is a two stage process; hard segments degrade first, soft polymer segments next. Also the thermal decomposition of adhesive tapes on the basis of acrylic rubber, acrylic-ethylene rubber, and acrylonitrile-butadiene rubber has been studied by TG-DTA [66]. Similarly, new block copolyurethane elastomers consisting of 1,2-ethylene bis (4-phenylisocyanate), poly(etherglycol) and ethylenediamine, were characterised by DSC, TG-DTA and FT-IR to determine the morphological structures and thermal stability. PTMG-based copolyurethane shows superior thermal stability due to its more cohesive hard domains and a better mutual-stabilisation effect between hard and soft segments [67]. TG-DTA results have also shown that the presence of SnO2 and/or γ-Fe2O3 could effectively inhibit the oxidative decomposition of residual Si-OH groups in an addition type silicone rubber and improve the thermal stability [68]. TG-DTA measurements have further been used to indicate that elastic polyaniline has excellent thermal stability [69]. The combination of TG-DTA with DSC, DMA and TMA is useful in monitoring and characterising the influence of artificial ageing on the chemical and physical properties of EPDM roofing materials. Significant changes in Tg, the coefficient of thermal expansion and the chemical composition were observed due to heat ageing [70]. Thermo-analytical techniques and infrared have been applied to investigate the effect on the molecular level of heat ageing, humidity and UV radiation over 8000 hours on several silicone-based sealants [71]. TG-DTA confirmed the susceptibility to loss of polydimethylsiloxane (PDMS) as well as other sealant components at elevated temperatures. Weathering of thermoplastic polyester elastomers (TPEE) has been studied by TG-DTA, FTIR, 1H nuclear magnetic resonance (NMR) and gel permeation chromatography (GPC) [72]. Ether parts of the soft segments in the polymer are degraded selectively and ester bonds are formed. There was a clear difference between tendencies of degradation by outdoor exposure and accelerated weathering. A TG-DTA study of the thermochemical processes occurring at vulcanisation temperatures with N-oxydiethylene-2-benzthiazyl sulphenamide and N-cyclohexyl-2-benzthiazyl sulphenamide and their mixtures with sulphur showed the formation of high molecular weight polysulphides [73]. The influence of metallic oxides (Fe2O3, SnO2) on hot air ageing of one-pack room temperature vulcanised fluorosilicone rubber has been studied by means of TG-DTA [74, 75]. TG-DTA and TG were both applied to study the thermal characteristics of room temperature vulcanised silicone rubber [76]. TG-DTA and DSC are suitable for product quality control as exemplified by OIT measurements for polyethylene (PE) and quantitative analysis of the rubber phase in ABS and of a polymer/ softener/soot/mineral filler mixture [77].

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Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… Khan [78] has examined the thermal response of various multiphase engineering materials (such as modified PVC/EPDM composites) in relation to the integrity and morphology of the interfacial junction. TG-DTA data displayed enhanced thermal stability by the EPDM insulation in the advanced viscoelastic composite that compares well with competitor advanced silicone composites. Riga and co-workers [79] have examined the oxidative behaviour of commercial engineering plastics, polyolefins and elastomers by simultaneous TG-DTA and pressure differential scanning calorimetry (PDSC). Riga [80] has also reported illustrative case studies showing how PDSC, modulated differential scanning calorimetry (MDSC), TMA, TG-DTA and DETA are used to select polymers, elastomeric automotive seals and lubricant additives. Elastomeric parts must be flexible and impervious to hostile engine environments, including hot-stressed motor oils possibly containing additives. Thermal analysis can be used to evaluate seals for composition (TGA), mechanical strength and glass transition temperature (DMA and/or TMA) and oxidative stability (PDSC and TG-DTA) [81]. Yang and co-workers [82] have reported a mechanistic study carried out based on the simultaneous measurement of enthalpy and weight loss by TG-DTA on tyre rubbers: NR, SBR and BR. The hyphenated technique allows to attribute the observed exothermic peak to chemical reactions and the endothermic peak to evaporation of pyrolysis products. The evaluation of combustion performance and safety of Mg/polytetrafluroethylene (PTFE) pyrotechnic compositions by means of TG-DTA has been reported by Miyake and co-workers [83]. Similar work on different pyrotechnic systems containing a chlorinated rubber binder has made use of TG-DSC and TG-DTA-MS [59-62]. The hyphenated technique TG-DTA (DSC) has recently been reviewed [56].

1.3.6 Thermogravimetry - Differential Thermal Analysis - Mass Spectrometry TG/DTG-DTA-MS equipment has been available since 1979 [11]. As mentioned before, simultaneous TG-DTA-MS was used to investigate the exothermic reactions that take place at approximately 300 °C between Alloprene binder and NaNO3 and Sr(NO3)2 in various pyrotechnic compositions, including the effect of titanium [85]. For both binary nitrate - Alloprene mixtures, reaction is initiated by Alloprene decomposition; another characteristic reaction is that of HCl (from the chlorinated rubber decomposition) with the nitrate to yield the metal chloride. Manley [32] examined a cured phenolic formaldehyde (PF) resin by means of TG-DTAMS observing a lower sensitivity of TG relative to DTA. However, new TGA instrumental developments have been reported since. The TG curve shows loss of phenol (MS evidence: m/z 94); DTA observes water (MS: m/z 18), ammonia (MS: m/z 17) and formaldehyde

33

Spectroscopy of Rubbers and Rubbery Materials (MS: m/z 29), indicating disrupture of cross-links greatly effecting the mechanical properties of PF moulded compounds. The MS traces show catastrophic deterioration of PF resins at 200 °C. The DTA trace also signals a change around 200 °C. DTA is thus a useful indicator of temperatures at which engineering properties may change but MS shows clearly why these changes occur.

1.3.7 Thermogravimetry - Gas Chromatography - Mass Spectrometry Volatile additives for vulcanised or unvulcanised rubbers can be accurately identified by TG or by controlled heating of a test sample in a sealed vial equipped with an overhead collecting headspace, transferring the heated volatile substances to a chromatographic column and analysing the separated volatile components emerging from the chromatograph column by various selective analytical detectors. Several illustrative examples were mentioned before. Although it is possible to obtain separate responses for several molecular species by TGMS, pyrolysis of samples such as polymers may give many products. It is an advantage if these can be separated between the thermal analysis unit and the mass spectrometer by gas chromatography. Gorman [175] has proposed a controlled thermal desorption and concentration method (essentially headspace) for separating volatile additives from vulcanisable rubber, in a TD-GC-MS configuration without the need for prior sample preparation such as milling, extraction or pyrolysis. Meuzelaar and co-workers [150] used high pressure TG-GC-MS to simulate solvent-free thermal and catalytic liquefaction reactions for coprocessing of waste polymers (non-vulcanised SBR alone or mixtures of waste plastics composed of PE, PS and waste rubber tyres) with coal.

1.3.8 Thermal Desorption - Gas Chromatography - Fourier Transform Infrared Spectroscopy/Mass Spectrometry The coupling of thermal desorption and identification techniques constitutes a powerful means for the detailed characterisation of outgassing processes with many potential applications in the field of rubbers. Gorman [175] has described a thermal desorption process for separating volatile components such as accelerator fragments, antioxidants and other organic additives from 2-10 g (un)vulcanised rubber samples for the purpose of analysing the separated volatile compounds. The process comprises sealing the test sample in a glass vial containing a controlled atmosphere and overhead headspace, followed by heating for complete desorption. The volatilised gases are then swept through a capillary GC column and analysed by MS, FTIR and FID using column splitters.

34

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… Jansen and co-workers [86] have evaluated temperature-controlled outgassing processes of plastics and rubbers using both off-line and on-line TD-GC-FTIR-MS. Decomposition of polyesterurethanes by means of TG-Tenax off-line sampling followed by TD-GCFTIR-MS revealed CO2, H2O, tetrahydrofurane, cyclopentanone, dicarbonic acid, aliphatic diols and esters [86]. The same authors have also described the detection of polychlorinated biphenyls (PCB) in 2,4-dichlorobenzoylperoxide cured silicone rubbers after outgassing products of a rubber silicone part obtained after desorption for 10 minutes at 200 °C in the thermal desorption cold-trap and subsequent analysis by means of TDGC-MS. Using a mass range of 290-294 Da the MS can be used as a selective detector for these substances.

1.4 Future Prospects of Hyphenated Thermogravimetric Techniques in Elastomer Characterisation As has been seen, competitive analytical systems correlating thermal and chemical behaviour are TG-MS, TG-FTIR, TG-GC, TG-GC-MS, etc. This list of techniques may easily be extended with various pyrolysis methods. It may be noticed that in particular combinations of thermal analytical instrumentation with MS provide a powerful means of analysing and characterising polymeric materials, providing both qualitative data and giving quantitative applications. TG-MS is not the solution to all problems in the study of polymer structure and degradation; other techniques will continue to give valuable contributions. However, in many instances this technique is a very valuable first experiment in order to identify reaction products and to define the underlying chemistry. It is therefore expected that TG-MS will keep a key position in problem solving for polymer analysts. Clearly though, routine application of TG-MS requires skilled and dedicated operators. Reliable TG-MS equipment is now commercially available. A limit of about 500 Da (bench-top range) is likely to be set by interface design requirements. The current interest of the material in events at lower temperatures (with evolution of complex mixtures of components at much lower concentration at the onset of the reaction) requires the need for quantification of components at very low concentrations and demands high sensitivity, accuracy, and reproducibility of the analytical technique; at the same time, the identification capability of the method requires strengthening. Today’s challenge is no longer faster collection of mass spectral data over broader (m/z) ranges, but more rapid identification of co-evolved species. In this respect several approaches for improvement are possible, such as high resolution TG-MS experiments (allowing for better experimental resolution of closely spaced thermal events), soft

35

Spectroscopy of Rubbers and Rubbery Materials ionisation techniques (increasing the identification power of the mass spectrometric event) and principle component analysis (PCA) of TG-MS data (for more reliable interpretation). Indeed, Alders and co-workers [151] have suggested that high resolution thermogravimetry – electron impact/soft ionisation quadrupole mass spectrometry (HRTG-EI/SI QMS) extended with PCA is currently the most advanced design for TG-MS coupling. PCA analysis of TG-MS data has recently successfully been carried out by Tas and co-workers [176]. The future will probably see a renaissance of TG-ToFMS. A major challenge in TG/DTG-based analysis of elastomer vulcanisates is to demarcate oil/plasticiser and elastomer regions, which often show overlapping TG events. Most of these materials have volatilisation ranges rather than discrete volatilisation points because they are chemical blends of components of various molecular weights and volatilities. Deconvolution of the overlapping oil/plasticiser and oil/elastomer TG curves is expected to be feasible with HRTG-MS-PCA which would substitute the dated methods for graphical resolution of oil and polymer weight loss [177]. The increasing need to correlate thermal behaviour simultaneously with the underlying chemistry (outgassing, thermostability, degradation) and the accompanying physical phenomena (blooming, cracking, foaming, migration, change in colour) determines the usefulness of a variety of endoscopic, audiometric and magnetometric extensions. For example, VI-TG is a direct visualisation of the morphological and textural changes in the substrate during thermal processing. VI-TG-MS combines chemical and physical information with visual observations [26] and is expected to gain rapidly in popularity on account of its close observancy of chemical and physical phenomena and for documentation purposes. Method standardisation of multihyphenated techniques remains an important issue.

1.5 Summary The experimentalist disposes of a variety of powerful thermogravimetric tools for the characterisation and study of the properties of elastomeric materials. Interpretation of TG data is facilitated on the basis of spectroscopic and mass spectrometry information which ease interpretation and quantification. The study of polymeric materials by a combination of thermal analysis and chemical identification tools provides the chemist with a wealth of information. In the case of TG-FTIR and TG-MS molecular weight information is collected concerning the evolved gases which are responsible for the detected weight losses. The chemical information obtained allows separation of concurrent or overlapping reactions. New ways have been indicated to allow for a complete identification of co-evolving species using simultaneous experimentation. Alternatively, it is

36

Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… advantageous to separate the components in a gas chromatograph before analysis in the mass spectrometer.

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Spectroscopy of Rubbers and Rubbery Materials 84. C.Y. Park. E.H. Hwang and S.K. Min, Eylasutoma, 1998, 33, 3, 201. 85. P. Emmott, T.T. Griffiths and J. Queay, Proceedings of the 17th International Pyrotechnics Seminar, Beijing, China, 1991, 25. 86. J.A.J. Jansen, W.E. Haas, H.G.M. Neutkens and A.J.H. Leenen, Thermochimica Acta, 1988, 134, 307. 87. A.E. Pavlath, K.S. Gregorski and R. Young, Thermochimica Acta, 1985, 92, 383. 88. H.K. Yuen and G.W. Mappes, Thermochimica Acta, 1983, 70, 269. 89. N.S. Arnold, M-G. Kim, W.H. McClennen, J.P. Dworzanski and H.L.C. Meuzelaar, Proceedings of the Workshop Ion Mobility Spectrometry, Mescalero, NM, USA,1992, 11. 90. K. Liu, E. Jacab, W.H. McClennen and H.L.C. Meuzelaar, ACS Division of Fuel Chemistry, Preprints, 1993, 38, 823. 91. F. Paulik, J. Paulik and L. Erdey, Zeitschrift für Analytical Chemistry, 1958, 160, 241. 92. F. Paulik, J. Paulik and L. Erdey, inventors; no assignee; HU Patent 145, 369, 1955. 93. I. Berecz, S. Bohátka, G. Langer and G. Szöor, International Journal of Mass Spectrometry and Ion Physics, 1983, 47, 273. 94. J.P. Redfern, International Labmate, 1986, 11, 1, 19. 95. Technical Data Sheet STA 409, Netzsch Gerätebau, Selb, Germany. 96. E. Kaisersberger and E. Post, Thermochimica Acta, 1997, 295, 1-2, 75. 97. J.A.J. Jansen and W.E. Haas, Analytica Chimica Acta, 1987, 196, 69. 98. D.A.C. Compton, International Labmate, 1987, 12, 4, 37. 99. R.C. Wieboldt, G.E. Adams, S.R. Lowry and R.J. Rosenthal, American Laboratory, 1988, 20, 1, 70. 100. P.R. Solomon, M.A. Serio, R.M. Carangelo, R. Bassilakis, Z.Z. Yu, S. Charpenay and J. Whelan, Journal of Analytical Applied Pyrolysis, 1991, 19, 1.

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Characterisation of Elastomers Using (Multi) Hyphenated Thermogravimetric… 133. H.G. McAdie, Analytical Chemistry, 1967, 39, 543; 134. H.G. McAdie, Analytical Chemistry, 1972, 44, 640. 135. H.G. McAdie, Analytical Chemistry, 1974, 46, 1146. 136 S. Affolter and M. Schmid in Developments in Polymer Analysis and Characterisation, Rapra Technology Ltd., Shawbury, UK, 1999, Paper 5. 137. B. Courtault, Analusis, 1979, 7, 481. 138. A.K. Sircar, Rubber Chemistry and Technology, 1992, 65, 3, 503. 139. ASTM D297-93, Standard Test Methods for Rubber Products – Chemical Analysis, 1998. 140. ASTM E1131-98, Standard Test Methods for Compositional Analysis by Thermogravimetry, 1998. 141. D.W. Brazier, Rubber Chemistry and Technology, 1980, 53, 3, 437. 142. M. Wingfield in Fachtagung. Kopplungen der Instrumentellen Analytik (TA, IR, MS) für die Kunststoff- und Kautschukindustrie, Ed., H. Möhler, Süddeutsches Kunststoff Zentrum, Würzburg, 1991. 143. R.G. Beimer, ACS Division of Organic Coatings and Plastics Chemistry, Preprints, 1975, 35, 428. 144. G.J. Mol, Thermochimica Acta, 1974, 10, 259. 145. M. Simpson, P.M. Jacobs and F.R. Jones, Composites, 1991, 22, 2, 105. 146. A. Kettrup and K.H. Ohrbach, Proceedings of the 8th International Thermal Analysis Conference, Alfa, Bratislava, 1985, 2, 629. 147. G.A. Kleineberg, D.L. Geiger and W.T. Gormley, Makromolekulare Chemie, 1974, 175, 2, 483. 148. G.A. Kleineberg and D.L. Geiger, Proceedings of the 3rd International Confernce on Thermal Analysis, Davos, Switzerland, 1971, 325. 149. M. Negri and F. Alarcon - Lorca, Revue Generale des Caoutchoucs et Plastiques, 1994, 61, 637/8, 55.

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Spectroscopy of Rubbers and Rubbery Materials 150. K. Liu, E. Jakab, W. Zmierczak, J. Shabtai and H.L.C. Meuzelaar, ACS Division of Fuel Chemistry, Preprints, 1994, 39, 2, 576. 151. A. Alders, K.G.H. Raemaekers and J.C.J. Bart, Proceedings of the SPE ANTEC ’99, New York City, NY, USA, 1999, Volume 3, p.3761. 152. D.J. Johnson and D.A.C. Compton, American Laboratory, 1991, 23, 1, 37. 153. S.M. Dyszel, Analaytical Calorimetry, 1984, 5, 277. 154. J. Morelli, Journal of Analytical Applied Pyrolysis, 1990, 18, 1. 155. M. Maciejewski and A. Baiker, Thermochimica Acta, 1997, 295, 1-2, 95. 156. M. Maciejewski, C.A. Müller, R. Tschan, W.D. Emmerich and A. Baiker, Thermochimica Acta, 1997, 295, 167. 157. R. Schönherr, TGA-FTIR Atlas Elastomere, Verlag W.K. Schönherr, Burgdorf, Germany, 1996. 158. D.D. Jiang, Q. Yao, M.A. McKinney and C.A. Wilkie, Polymer Degradation and Stability, 1999, 63, 423. 159. B.J. McGrattan, Applied Spectroscopy, 1994, 48, 12, 1472. 160. R.G. Davidson, Mikrochimica Acta, 1988, 1, 301. 161. J.W. Washall and T.P. Wampler, Spectroscopy, 1991, 6, 38. 162. B. Bowley, E.J. Hutchinson, P. Gu, M. Zhang, W.P. Pan and C. Nguyen, Thermochimica Acta 1992, 20, 209. 163. E.G. Jones, D.L. Pedrick and I.J. Goldfarb, Polymer Engineering and Science, 1988, 28, 1046. 164. C. Chang and J.R. Tackett, Thermochimica Acta, 1991, 192, 181. 165. H.K. Yuen, G.W.Mappes and W.A. Grote, Thermochimica Acta, 1982, 52, 143. 166. K-H. Ohrbach and A. Kettrup, Polymer Degradation and Stability, 1985, 13, 99. 167. W. Holzapfel, in Fachtagung. Kopplungen der Instrumentellen Analytik (TA, IR, MS) für die Kunststoff- und Kautschukindustrie, Ed., H. Möhler, Süddeutsches Kunststoff Zentrum, Würzburg, 1991.

46

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48

2

Photoacoustic Fourier Transform Infrared Spectroscopy of Rubbers and Related Materials James R. Parker

2.1 Introduction Photoacoustic Fourier transform infrared spectroscopy (PA-FTIR) differs from most infrared techniques in that it is an emission rather than an absorption technique. It is a highly versatile technique requiring little sample preparation other than sizing the sample. It can be used for both qualitative and quantitative purposes and for surface and depth profiling analyses. It has an interesting history with the first crude photoacoustic spectrometer being built in the 1880s by Alexander Graham Bell, the inventor of the telephone. In this chapter the basic principles and typical uses of what is known as rapid scan PA-FTIR are presented. The analyses are those that can be carried out with any mid-level grade Fourier transform-infrared (FTIR) spectrometer and a photoacoustic detector.1

2.2 History of Photoacoustic Spectroscopy The photoacoustic effect was discovered by Alexander Graham Bell in 1880 [1]. It led to the development in 1881 of what Bell termed the spectrophone [2]. It had been found that solids, liquids, and gases when irradiated with modulated sunlight produced sounds of varying intensities depending upon the material. Bell realised that there was a relationship between the radiation absorbed and the sounds produced and that the technique would be useful in studying absorption spectra in the infrared. There appear to have been several versions of the spectrophone and the version described in a biography is a spectrophotometer utilising a prism to isolate wavelengths [3]. Others confirmed Bell’s results and postulated various theories to explain the results [4-8]. At this point it was clear that the photoacoustic effect was most pronounced with gases. It was also well understood that the sound produced was the result of pressure/volume changes of the

1

‘Permission for the publication herein of Sadtler Spectra has been granted, and all rights reserved, by BIO-RAD Laboratories, Sadtler Division.’

49

Spectroscopy of Rubbers and Rubbery Materials gas. However, the mechanism involving solids and liquids was not well understood. Practical instrumentation for the determination of gases began to be developed in the late 1930s and early 1940s by Veingerov, Pfund, and Luft [9-12]. This instrumentation using typical infrared sources and microphone detectors developed the capability of determining parts per million of gases. Such gas detectors are still available. If lasers are used as sources than the detection limits can be pushed to the parts per billion level [13,14]. Satisfactory results with solids awaited the development of nonresonant photoacoustic cells by Rosencwaig and Gersho [15-17]. It awaited the further development and utilisation of FTIR spectrometers. Such work began to be reported in the late 1970s [18-21]. The theoretical basis for the photoacoustic effect in condensed phases was begun by Parker [22]. A general theory was developed by Rosencwaig and Gersho (the RG theory) and expanded upon by others [23-27].

2.3 Theory of Photoacoustic Spectroscopy In photoacoustic FTIR spectroscopy the sample in a sealed chamber (as illustrated in Figure 2.1) absorbs modulated infrared radiation. The excited molecules give up that excess energy as heat through collisional processes. The heat diffuses to the surface of the sample where it creates a pressure wave (an acoustic

Figure 2.1 Schematic of photoacoustic (PA) detector

50

Photoacoustic Fouriers Transform Infrared Spectroscopy of Rubbers and Related Materials wave) which is then detected by a microphone. It is readily seen that any theory will have to be fairly complicated and involve a number of optical and thermal parameters. The RG theory uses the following parameters: D the thickness of the sample (cm) Do the optical absorption length is the distance the infrared radiation can travel through the sample (cm), equal to 1/β where β is the optical absorption coefficient Dt the thermal diffusion length is the distance the thermal wave can travel through the sample (cm) The book by Rosencwaig can be consulted for a thorough discussion of the theory [27]. To simplify the theory, samples are categorised depending on the relative magnitude of the following parameters: Optically transparent

Do > D, the optical absorption coefficient is small

Optically opaque

Do < D, the optical absorption coefficient is large

Thermally thin

Dt > D, thermal conductivity is large

Thermally thick

Dt < D, thermal conductivity is small

The typical polymer or rubber sample would be classified as optically transparent or opaque and thermally thick except possibly for the strongest bands. In this case the signal intensity would be proportional to the product of the optical absorption coefficient (β) and the thermal diffusion length and show a – 3/2 dependence on the modulation frequency (ω). The angular modulation frequency is a product of the interferometer mirror velocity and the wavenumber: ω = 2π (wavenumber, cm-1) (OPD velocity, cm/s)

(2.1)

where OPD is the optical path difference = 4π (wavenumber, cm-1)(mirror velocity, cm/s)

(2.2)

for a Michelson interferometer Sometimes the mirror velocity must be calculated from what the instrument manufacturer specifies as the He-Ne laser modulation frequency. Mirror velocity, cm/s = [(laser modulation frequency, KHz )(1000)] ÷ [(2)(15798)]

(2.3)

51

Spectroscopy of Rubbers and Rubbery Materials The thermal diffusion depth is also related to the modulation frequency: Dt = (2α/ω)1/2

(2.4)

where α is the thermal diffusivity (cm2/s) of the sample. The thermal diffusivity is calculated from other thermal parameters and the density. α(cm2/s) = (thermal conductivy,cal/cm-s-°C) ÷ [(density, g/cm3) (specific heat, cal/g-°C)] (2.5) Since only the radiation absorbed within the thermal diffusion depth gives rise to the signal, the depth being sampled can be varied by changing the modulation frequency. The sample spectrum is ratioed against that of carbon black or a polymer or rubber sample highly filled with carbon black. This type of sample is classified as optically opaque and thermally thin. Here the signal intensity is independent of the optical absorption coefficient, photoacoustic saturation has taken place, and shows a –1 power dependence on the modulation frequency. This single beam spectrum resembles the one obtained when a background spectrum is obtained using a deuterated triglycine sulfate (DTGS) detector, i.e., a blackbody emission curve. When the sample spectrum is ratioed against the carbon black reference spectrum the resulting spectrum has a signal intensity with a –1/2 power dependence on the modulation frequency (-3/2 ÷ -1 = -1/2). Chalmers and co-workers, used the spectrum of polyvinyl chloride to test this dependence and found good agreement for all the bands tested (963, 1100, 1331, and 1431 cm-1) except for the strongest band (1250 cm-1) where saturation may have occurred [28]. The signal increased with decreasing wavenumber in the manner predicted.

2.4 Instrumentation for PA-FTIR Analyses A wide variety of FTIR instruments have been used to obtain PA-FTIR spectra. A cursory examination of the literature reveals some of the instruments used: JEOL JIR-5500; IBM IR-95, IR-98, 9195; Bruker IFS 66, IFS 88, IFS 113V; Perkin Elmer 1750, 1760-X, 1800; Bomen DA 3.02; Nicolet 20 SX, 20 DXB, SX-170, 7199, 740, 800; Mattson Cygnus 100, Sirius 100, Polaris; Digilab FTS-10, FTS-10M, FTS-11, FTS-15, FTS-20, FTS-20E, FTS-60, FTS-65, FTS-6000; Laser Precision Analytical (Analect) RFX-75 and Analect FX-260. The Laser Precision Analytical Instrument required extensive modification. The list is only intended to indicate the wide variety of different instruments which have been used and may not be complete. The normal adjustments for maximum performance should be made, beam splitter alignment, etc. The source aperture, if any, should be opened fully for maximum source intensity. The gain control and other adjustments for best operation at each of the different mirror velocities to be used should be made. The

52

Photoacoustic Fouriers Transform Infrared Spectroscopy of Rubbers and Related Materials instrument should be set on a solid table to minimise vibration and purged with dry air or nitrogen. Prior to use of the PA detector the bench should be initialised with the DTGS detector to establish that it is operating normally. Some of the commercial photoacoustic detectors listed in the literature are: Ames; EDT OAS 401; LPA FX-400; Digilab; PAR 6003; Gilford; Nicolet; MTEC 100, 200, and 300. MTEC appears to be the only commercial producer at the moment. It is not clear that the results from all these detectors are equally reliable and comparable although reviewers tend to treat them as such. Figure 2.1 shows a general schematic of a photoacoustic detector. The MTEC detector has a BNC connector (a metal device used to connect a coaxial cable to something else) which is usually interfaced with the external detector port of the FTIR. A second cable is connected to the MTEC power supply which is used to power the photoacoustic detector and to remotely control the gain of the preamplifier in the detector. It is this gain control which is adjusted when obtaining spectra of samples for the appropriate signal level. Usually the detector is purged with helium because of its good thermal properties and the instrument is operated with a low mirror velocity, e.g., 0.1 cm/s, and a resolution of 8 cm-1 over a range of 450-4000 cm-1 with 32 scans. After the detector has been aligned and is ready for sample data collection, a background spectrum of the carbon filled polymer standard supplied by MTEC is obtained. Alternatively carbon black powder, pelletised carbon black, carbon filled rubber, pyrolytic graphite, and other totally infrared absorbing materials have been used. A typical background spectrum is shown in Figure 2.2 with and without the use of a helium purge.

Figure 2.2 Single beam PA-FTIR spectra of MTEC carbon-filled polymer

53

Spectroscopy of Rubbers and Rubbery Materials Helium gave a 2.3 X improvement in sensitivity which is similar to the 2.5 X factor found by Wong [29]. Early workers attributed the high sensitivity obtained with helium to be due to its high thermal conductivity and low heat capacity [29, 30]. The other purge gas which has sometimes been used is xenon, a highly polarisable gas. Its use enhances the intensities of groups perpendicular to the surface and has been useful in studying the orientation of hydroxyl groups and silane coupling agents on silica surfaces [31, 32]. The sample can be of any shape or form as long as it fits into the sample holder which is 10 mm in diameter. The thickness can be up to 6 mm. The standard large sample cup is 3 mm in depth. There may be some optimum distance between the sample and the top of the cup for optimum generation of a signal in the gas layer. Carter and Wright tested the hypothesis that the optimum distance would be the thickness of the gaseous boundary layer (BLg) [33]. It is the expansion and contraction of this layer that generates an acoustic signal which is detected by a microphone. The thickness of the boundary layer is equal to 2πDg where Dg is the thermal diffusion length of the gas. BLg = 2πDg = (2α) ÷ [(4π)(wavenumber, cm-1 )(mirror velocity, cm/s)]

(2.6)

In this equation, α is the thermal diffusivity of the gas which is 1.51 cm2/s for helium. Therefore, at a mirror velocity of 0.08 cm/s and a wavenumber of 500 cm-1, the thermal diffusion length is 0.78 mm and the boundary layer thickness is 4.9 mm. Carter and Wright concluded that the optimum distance was somewhat less than that of the boundary layer thickness and that other factors were involved. McClelland and co-workers, have suggested that a distance of 1mm or the thermal diffusion length of the gas be used [34]. The sample used for analysis can be in a wide variety of forms. These can include irregular shaped solids, pellets, powders, films, fibres, or liquids and need not be transparent. The primary restrictions be that the sample fits in the sample holder and that it not degas since the spectra of gases are stronger than those of solids. For samples containing moisture this might require oven drying, using a desiccant beneath the sample in the sample holder, or simply allowing the sample to equilibrate in the instrument atmosphere prior to insertion in the photoacoustic detector. It does not mean that hydrates cannot be analysed. Powders give spectra similar to those of obtained for solid samples as shown for polystyrene in Figure 2.3. In general increasing surface area increases band intensity. For example, the C-H band intensity of hydrogenated diamond powders correlated with the Brunauer, Emmet and Teller (BET) surface area over a range of 0.2 to 22 m2/g [35]. For some types of powders other factors also appear to be involved with signal enhancement. Yang and Fateley reported that some type of extra signal enhancement occurred for quartz powder when the particle diameter became smaller than the thermal diffusion length for a given peak [36]. A more general theory is that in addition to the thermal signal predicted by the RG theory there is a pressure signal caused by the interstitial gas which depends upon the porosity of the powder [37, 38]. Pandurangi and Seehra showed that the change in signal

54

Photoacoustic Fouriers Transform Infrared Spectroscopy of Rubbers and Related Materials

Figure 2.3 Effect of sample shape on PA-FTIR spectrum of polystyrene intensity of silica powders of varying porosity (ε) has an (ε)/(1-ε) dependence for strong bands and an e dependence for weaker bands in accordance with this theory [39]. Hövel, Grosse, and Theiss using a somewhat different theoretical approach also concluded that the expansion of the interstitial gas in porous samples is responsible for a significant portion of the photoacoustic signal [40]. Perhaps a final consideration when obtaining of PA-FTIR spectra is the FTIR computer software. Most IR experiments involve measuring transmitted radiation while the photoacoustic spectroscopy (PAS) experiment involves measuring emitted radiation. On some newer instruments the software may already be set up to take this into account. On others a bit of trickery is involved. The Mattson Cygnus 25 (Rev. 8) with First DOS based software is an example of one of the latter. When obtaining a PA-FTIR spectrum the instrument parameters in the software are set up to obtain a transmittance (.ras file) spectrum. The resulting sample spectrum is obtained and ratioed against the carbon black or other reference spectrum. The resulting .ras file must now be converted to an absorbance file (.abs) not through spectral conversion routines but by changing the data file extension to .abs. In First software this requires writing a macro or batch file routine for this purpose. The Y-axis may then be relabelled relative intensity and the scale may be changed if that aids in the utilisation of some desired software routine (library searching, etc). In any case the Y-axis is treated as being proportional to concentration. Typical PAS spectra for natural rubber (NR) and silica are shown in Figures 2.4 and 2.5.

55


Figure 2.4 PA-FTIR spectra of natural rubber (NR) Upper spectrum reproduced with permission from Bio-Rad Laboratories, Sadtler Division

Figure 2.5 PA-FTIR spectra of amorphous silica Upper spectrum reproduced with permission from Bio-Rad Laboratories, Sadtler Division

56

Photoacoustic Fouriers Transform Infrared Spectroscopy of Rubbers and Related Materials In the spectrum of natural rubber, the bands of particular importance are the C=C stretching band at 16663 cm-1, the C-H bending band of the methyl group at 1378 cm-1, and the C-H out-of-plane bending band from the cis unsaturation configuration at 837 cm-1. The spectrum of silica shows a broad band at 1100 cm-1 from the Si-O-Si antisymmetric stretching vibration and a band at 803 cm-1 from the Si-O-Si symmetric stretching vibration. There is a very broad band centred around 3376 cm-1 which represents the O-H stretching band of bound and unbound water and the Si-OH stretching band of silanols involved in various degrees of hydrogen bonding to water and other silanol groups. The distinct shoulder band at 3735 cm-1 represents the silanol groups which are least hydrogen bonded. There is a small Si-O-Si combination band at 1872 cm-1 and a water O-H bending bond at 1636 cm-1 which is on top of some type of Si-O overtone or combination band. The band at 945 cm-1 decreases in size as drying takes place and the major part of it is believed to represent the hydrogen bonding of water to silanols. These spectra were converted from .ras to .abs files with no change in the Y-axis scale and then used to search Sadtler libraries of commercial material using Mattson First software with a correlation coefficient search algorithm which takes into account baseline drift and differences in scaling. As can be seen the sample spectrum and the library spectrum are usually quite similar. The library spectra shown here were the best hits in each of the searches with a search correlation coefficient of >0.9 with 1.0 representing a perfect match.

2.5 Analysis of Carbon-Filled Rubbers Carbon-filled rubber samples are quite common. Many different methods for the IR analysis of these materials have been used including use of microtomed sections, attenuated total reflectance (ATR), specular reflectance, potassium bromide pellets, solubilisation with o-dichlorobenzene, and pyrolysis. None of these methods is totally satisfactory and there is a constant search for others. Problems arise because carbon black is such a strong absorber of IR radiation, scatters radiation, raises the refractive index of the sample, and is used in high enough concentrations to simply reduce the amount of sample available for examination. Photoacoustic spectroscopy is another method which has been examined since it is relatively simple to use and has a wider wavelength range than most ATR methods and some MCT detectors used in conjunction with other methods. Teramae examined nitrile-butadiene rubbers (NBR) containing up to 29 wt.% carbon black [41, 42] and Carter and co-workers, examined natural rubbers (NR) containing up to 25 wt.% carbon black [43]. Waddell and Parker give an example where the presence of NR can be detected in a sample containing 38.4 wt.% carbon black [44]. Figure 2.6 shows examples of PA-FTIR spectra of styrene-butadiene rubbers (SBR) with 0, 10, 20, and 30 wt.% carbon black.

57


Figure 2.6 Effect of carbon black on PA-FTIR spectrum of styrene-butadiene rubber (SBR)

The maximum concentration of carbon black that can be tolerated has not been established. Several ways of improving detectability have been proposed. The simplest of these is the one suggested by McClelland and involves obtaining a single beam spectrum of the sample and creating a background spectrum through the use of smoothing routines [45]. He reports that the resulting sample/background ratioed spectrum looks something like a derivative spectrum. Another possible way of creating a background spectrum is to calculate one from the sample spectrum as suggested by Coates [46]. He has illustrated this technique by using it to correct the severe curvature of a transmission spectrum of a carbon-filled ethylene-vinyl acetate (EVA) copolymer. Finally, there is the technique of linearising photoacoustic spectra. This involves recording photoacoustic spectra at two different mirror velocities and requires the use of some special software which makes use of phase information. Carter has used this technique to show the improvements available in spectra of formulated NR containing up to 25 wt.% carbon black [47]. PAS-FTIR spectra have been used to find out the interaction of chlorosulphonated polyethylene (CSM) and carbon black N110 [48]. A number of bands in the 1800 cm-11680 cm-1 region in the spectrum of N110 (Figure 2.7) confirm the presence of different carbonyl functionalities, which may include carboxyl group, lactone and quinone. The band at 1651 cm-1 is characteristic of aromatic double bonds in the carbon black. The

58


Figure 2.7 FTIR-photoacoustic spectra of pure CSM, b) N110 carbon black, c) CSM/ N110 carbon black compound

59

Spectroscopy of Rubbers and Rubbery Materials spectrum of the carbon black filled CSM compound retains all the characteristics bands due to the carbon black and CSM. However, there occur several additional bonds in the 1100-1000 cm-1 region, which indicate the formation of an ether. The –OH on carbon black reacts with –SO2Cl group of the CSM, which on heating eliminates SO2, forming an ether group, C-O-C, indicating polymer filler interaction, according to following equation:

(2.7)

2.6 Quantitative Analysis of Polymers There is not an extensive body of literature on the quantitative PA-FTIR analysis of rubbers and polymers. Early in the development of the technique there were indications that it might be possible. Many of the early studies involved the use of vinyl acetate copolymers. Vinyl chloride-vinyl acetate copolymers were briefly examined by Chalmers and co-workers, in 1981 by PA-FTIR with indications that quantitation might be possible [28]. They were also studied by Kirkbright and Menon in 1982 using near-IR photoacoustic spectroscopy [49]. In this case the intensity of the second overtone of the carbonyl band (4651 cm-1) was ratioed to the intensity of the first overtone of the C-H stretching vibration (5714 cm-1) and correlated with the vinyl acetate content. A chemometric PA-FTIR approach was used in 1991 by McClelland and co-workers to determine the percentage vinyl acetate in EVA over a range of 10%-50% [50]. They used three spectral ranges to represent compositional variations: (1) 2750-3120 cm-1, C-H stretching (2) 1630-1900 cm-1, C=O stretching and (3) 578-1490 cm-1, molecular vibration fingerprint region. In addition the band area in the 2750-3120 cm-1 region was used to correct for experimental variabilities. Cross-correlation analysis gave standard errors of prediction of 0.64%. In 1996 Parker and Waddell used two approaches for the PA-FTIR determination of vinyl acetate in EVA [51]. The intensity ratio of the 610 cm-1 acetate group bending band to the 722 cm-1 methylene group in-phase rocking vibration was used to determine 10%-50% vinyl acetate. For levels below 10% (3%-9%), the intensity of the carbonyl stretching band at 1740 cm-1, was ratioed to the intensity of the methylene bending vibrational band at 1466 cm-1. Least squares linear regression analyses were used in each case with correlation coefficients of 0.99 and 0.97 being obtained, respectively.

60

Photoacoustic Fouriers Transform Infrared Spectroscopy of Rubbers and Related Materials % Vinyl acetate = 38.8 x (610/722 Band Ratio) + 3.42

(2.8)

correlation coefficient = 0.992 % Vinyl acetate = 9.25 x (1740/1466 Band Ratio) – 3.07

(2.9)

correlation coefficient = 0.973 An example of the PA-FTIR analysis of a rubber is the determination of acrylonitrile (ACN) in NBR. A PA-FTIR spectrum of a NBR is shown in Figure 2.8. Parker and Waddell investigated three methods [51]. These involved determining the %ACN as a function of: (1) the intensity ratio of the 2238 cm-1 nitrile stretching band to the 1443 cm-1 methylene deformation or scissoring band (2) the intensity ratio of the 2238 cm-1 nitrile stretching band to the 969 cm-1 C-H outof-plane bending band of the trans-butadiene double bond and (3) the absolute intensity of the 2238 cm-1 nitrile stretching band.

Figure 2.8 PA-FTIR spectrum of nitrile-butadiene rubber (NBR) Reprinted with label modifications from reference [49] with permission from Technomic Publishing Co., Inc., © Copyright 1999

61

Spectroscopy of Rubbers and Rubbery Materials Least squares linear regression analyses of the data over the range of 20%-50% ACN gave the following equations: %ACN = 60.5 x (2238/1443 Band Ratio) – 0.271

(2.10)

correlation coefficient = 0.9995 %ACN = 79.74 x (2238/969 Band Ratio) + 5.08

(2.11)

correlation coefficient = 0.999 %ACN = 4.84 x (2238 Band Intensity) + 1.99

(2.12)

correlation coefficient = 0.97 Another example of the PA-FTIR analysis of a rubber is the determination of % styrene in SBR. A typical PA-FTIR spectrum of an emulsion SBR is shown in Figure 2.9. Traditional non-PA methods might involve measuring the 700 cm-1 styrene band in carbon disulfide solutions as carried out by Binder [52], transmission measurements on films

Figure 2.9 PA-FTIR spectrum of emulsion styrene-butadiene rubber (SBR) Reprinted with label modifications from reference [49] with permission from Technomic Publishing Co., Inc., © Copyright 1999

62

Photoacoustic Fouriers Transform Infrared Spectroscopy of Rubbers and Related Materials deposited from solution as done by Panyszak and Kovar where the 970 cm-1 butadiene band was ratioed to the 755 cm-1 styrene band [53], or an ATR method like that of Sloan and Clements where the ratio of the 698 cm-1 styrene band to the 967 cm-1 trans-1,4butadiene band was used [54]. Parker and Waddell studied the PA-FTIR analyses of cold process emulsion SBR [51]. The styrene content was determined by ratioing the intensity of the 698 cm-1 styrene band to the 968 cm-1 trans-1,4-butadiene band as a function of % styrene. The results can be expressed by a least squares linear regression equation over the range of 5%-56% styrene in SBR. % Styrene = 53.6 x (698/968 Band Ratio) – 13.54

(2.13)

correlation coefficient = 0.997 An alternative method was also studied. This involved ratioing the intensity of the 698 cm-1 styrene band to the intensity of the 2921 cm-1 C-H stretching vibration. Since oils and other additives would interfere with this approach they were extracted with acetone. Vacuum oven drying was then necessary to remove all traces of acetone prior to PA analysis. Otherwise the PA spectrum would be that of acetone rather than that of the rubber since the gas phase spectrum of the acetone would overwhelm that of the solid phase rubber. This technique allowed both solution and emulsion SBR to be analysed by a common method. The results can be expressed by a least squares linear regression equation over the range of 5%-40% styrene in SBR. % Styrene = 53.0 x (698/2921 Band Ratio) – 4.52

(2.14)

correlation coefficient = 0.998 Another example of rubber analysis is the determination of vinyl-butadiene in SBR. A PA-FTIR spectrum of a high-vinyl content SBR is shown in Figure 2.10. This determination by mid-IR methods invariably involves the use of the 910 cm-1 band due to the wagging motion of the pendent vinyl group of 1,2-butadiene. These methods have been reviewed by Silas, Yates and Thornton [55]. In a PA-FTIR method Parker and Waddell also used the intensity of the 910 cm-1 butadiene band and ratioed it to the intensity of the 1450 cm-1 C-H bending band as a function of the % vinyl-butadiene [51]. The results can be expressed by a least squares linear regression equation over a range of 10%-60% vinyl-butadiene. % Vinyl-butadiene = 28.4 x (910/1450 Band Ratio) – 8.16

(2.15)

correlation coefficient = 0.981

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Figure 2.10 PA-FTIR spectrum of high vinyl styrene-butadiene rubber (SBR) Reprinted with label modifications from reference [49] with permission from Technomic Publishing Co., Inc., © Copyright 1999

A final quantitative example is that of determining the % ethylene in ethylene-propylene copolymers (EPM). The International Institute of Synthetic Rubber Producers’ (IIRP) Technical/Operating Committee and the American Society for Testing and Materials (ASTM) cooperated to produce an updated standard method [56]. They also produced a set of standards with the ethylene content established by 13C NMR. Those standards were utilised by Parker and Waddell to study the photoacoustic determination of % ethylene [57]. Three of the four infrared band ratios specified by ASTM were measured, 1378/1462, 1378/722, and 1156/722. The fourth one, 1156/4255, was not used since it involves measurements in the near-IR region. Baselines were drawn by the valley-tovalley technique as illustrated in Figure 2.11. The 1378 cm-1 band is from the CH3 symmetric bending and the 1156 cm-1 band is a complex skeletal vibration involving the CH3 branch of propylene. The 722 cm-1 band represents the CH2 rock and the 1462 cm-1 band is a combination of the CH2 scissor and the asymmetric CH3 bend. In the photoacoustic spectra the 1378 and 1462 bands are strong while the 1154 and 722 cm-1 bands are weak. Least squares linear regression

64


Figure 2.11 Baseline constructions for PA-FTIR analysis of ethylene-propylene copolymers (EPM)

analysis of the three band ratios versus % ethylene shows the best correlation coefficient (0.9926) for the 1378/1462 band ratio, the next best (0.9670) for the 1378/722 band ratio, and the worst (0.9566) for the 1156/722 band ratio. These are as would be expected from the relative intensities of the bands involved. The corresponding equations for these correlations are as follows: % Ethylene = (121.35) – (100.22)(1378/1462 Band Ratio)

(2.16)

% Ethylene = (80.46) – (8.573)(1378/722 Band Ratio)

(2.17)

% Ethylene = (76.97) – (61.88)(1156/722 Band Ratio)

(2.18)

Quantitation of PA-FTIR data appears to be generally applicable in the analysis of rubbers and polymers. The chances of success are best if the strongest, possibly saturated, infrared bands are avoided. The quality of the results are improved if band ratioing is used to compensate for experimental variables.

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2.7 Surface Analysis and Depth Profiling A discussion of surface analysis requires a review of the depth being sampled during PAFTIR spectroscopy. The depth being sampled during PA-FTIR analyses of rubbers is the thermal diffusion depth (Dt). This is a function of the thermal diffusivity of the sample, the wavenumber, and the mirror velocity. Dt = (2α/ω)1/2 where ω = 4π(wavenumber)(mirror velocity)

(2.19)

Dt = [(2α, cm2/sec)] ÷ [(4π)(wavenumber, cm-1)(mirror velocity, cm/sec)] For illustrative purposes a thermal diffusivity of 1.3 x 10-3 cm2/s is often used as being typical of rubber and polymers. Some values from the literature for various materials are given in Table 2.1. Using the value of 1.3 x 10-3 it can be calculated that a depth of 3 to 11 μm is being sampled at 2000 cm-1 as indicated in Table 2.2. This is an order of magnitude greater than that sampled by ATR techniques. One type of sample to be considered is one which has a bloom or similar surface layer. This was an area first explored by Carter and co-workers [60]. They reported finding a bloom of dimorpholinyl thione and zinc stearate formed from ingredients used during the vulcanisation of the NR sample. As the mirror velocity was increased from 0.316 cm/s to 1.216 cm/s the bands associated with the bloom increased. Figure 2.12 shows the PAFTIR spectrum of a zinc stearate bloom on an amber coloured silica-filled SBR.

Table 2.1 Thermal diffusivities of rubbers and polymers Thermal Diffusivity (cm2/s x 10-3)

Reference

Polyethylene (PE), high density

2.3

[58]

Polyethylene, low density

1.4

[58]

Polypropylene (PP)

1.2

[58]

Polyethylene terephthalate

1.1

[58]

Polystyrene

1.1

[58]

Polytetrafluoroethylene

0.95

[58]

SBR

1.23

[59]

BR

1.0

[59]

Polyisoprene

0.86

[59]

Polychloroprene (Neoprene)

0.68

[59]

Rubber/Polymer

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Table 2.2 Thermal diffusion depth as a function of mirror velocity (for thermal diffusivity = 1.3 x 10-3) Thermal Diffusion Depth (μm)

Mirror Velocity

LMF

(cm/s)

(kHz)

4000 cm-1

2000 cm-1

500 cm-1

0.084

2.7

7.8

11.1

22.2

0.115

3.6

6.7

9.5

19.0

0.180

5.7

5.4

7.6

15.2

0.316

10

4.0

5.7

11.4

0.633

20

2.9

4. 0

8.4

1.264

40

2. 0

2.9

5.7

LMF: laser modulation frequency

Figure 2.12 PA-FTIR spectrum of bloom on rubber

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Spectroscopy of Rubbers and Rubbery Materials The 1539 cm-1 zinc stearate band can be ratioed to the 966 cm-1 SBR band and plotted as a function of mirror velocity (depth) to show that it is increasing in concentration as the surface is approached. This is shown in Figure 2.13. Another type of sample for surface analysis is one where some kind of reaction has taken place on the surface. A common example of this is the appearance of carbonyl bands from the surface oxidation of PE. A spectrum of ozone treated PE film is shown in Figure 2.14. The intensity of the 1738 cm-1 carbonyl band can be ratioed against the 1466 cm-1 PE band and plotted as a function of mirror velocity (depth) to see that the oxidation decreases with depth (Figure 2.15). Delprat and Gardette studied the photo-oxidation of PP with and without titanium oxide being present [61]. The intensity of the carbonyl band at 1713 cm-1 was followed using a mirror velocity of 0.16 cm/s. As they point out roughened surfaces do not lend themselves well to ATR techniques and microtoming may damage surface layers. Gonon, Vasseur, and Gardette studied the photo-oxidation of styrene-isoprene copolymers by monitoring the ratio of the 1718 cm-1 carbonyl band to the 1600 cm-1 aromatic ring band. Mirror velocities from 0.0158 to 0.6329 cm/s were used [62]. A somewhat different type of

Figure 2.13 Zinc stearate bloom concentration as a function of depth

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Figure 2.14 PA-FTIR spectrum of ozone treated polyethylene

Figure 2.15 Polyethylene oxidation as a function of depth

69

Spectroscopy of Rubbers and Rubbery Materials profiling was carried out by de Oliveira, Pessoa, Vargas and Galembeck [63]. They used near-infrared photoacoustic spectroscopy to determine the depth profile of –CH3, =CH2, and –OH groups in low density polyethylene over the 11-56 μm range [63]. They concluded that the surface is richer in –CH3 and =CH2 groups than the bulk and play a role in poor adhesion to polyethylene surfaces. Laminates have been extensively studied by PA-FTIR and are the subject of extensive investigations by the new much more complex step-scan PA-FTIR techniques. Kapton film is often used to illustrate the analysis of this type of sample. Kapton 200FN919 film is a DuPont product which consists of a 25 μm layer of polyimide laminated between 12.7 μm layers of Teflon. If a high mirror velocity is used then a spectrum of only the Teflon layer is obtained. On the other hand if a slow mirror velocity is used a composite spectrum of the Teflon and polyimide layer is obtained. The difference spectrum (low mirror velocity spectrum minus high mirror velocity spectrum) represents that of the polyimide layer. As illustrated in Figure 2.16 the difference spectrum closely resembles the PA FTIR spectrum of polyimide except for the regions where the strong C-F bands of the Teflon are present.

Figure 2.16 PA-FTIR spectra of Kapton film

70


2.8 Determination of Orientation Function of Polymeric Materials Photoacoustic spectroscopy with polarised light offers some potential for a new kind of investigation. For polarisation of IR radiation a KRS-5 wire grid polariser is inserted immediately in front of the cell. Without any special preparation polarised PAS can be directly applied to polymer granules, blocks, strings, sheets and plates as received from extrusion or injection moulding processing. In this way the preferred orientation given by the machine axis or stretching direction can be correlated with the polymer structure and processing conditions. Klaus-Jochen Eichhorn and co-workers have investigated the orientation of polymer components in different blends of thermoplastic polyurethane (TPU) with common and modified polypropylene (PP) after processing and in some cases after subsequent mechanical stretching [64]. TPU and PP form an immiscible blend but with maleic anhydride grafted PP (MSA), properties are improved. Polarised PAS is used to study the orientational behavior of TPU in the reactive compatibilised blend. The orientation of the polymers and their individual components in the blends has been investigated by means of the TPU bands between 1650-1800 cm-1 (carbonyl stretching vibrations), at 1316 cm-1 (C-N), 1115, 1081 cm-1 (C-O-C), respectively, and using the PP band at 1379 cm-1 (methyl deformation vibration). The direction of the dipole transition moments of the carbonyl, NH stretching and methyl deformation vibration is located in the perpendicular direction (⊥) relative to polymer chains axis and the direction of the others is located in the parallel direction (||), details are given in reference [64]. Hence polarised FTIR-PAS spectroscopy is suitable for detecting orientation effects in polymers and polymer blends, during processing like extrusion and injection molding. With polarised PAS, the difference in the orientation behavior of the hard and soft segments in a thermoplastic poly(ether urethane) can be determined [64]. Grafting of PP with maleic anhydride gives rise to a better orientation of the TPU hard segments in the blends.

2.9 Conclusion PA-FTIR has been shown to be a very versatile IR technique. All the usual types of studies can be carried out including qualitative and quantitative, surface and bulk analyses. The spectra can be searched, subtracted, annotated and otherwise manipulated although some software may require a change of the data file extension. The lack of necessity for any sample preparation offers many advantages which include allowing the uncontaminated sample to be analysed by other techniques. No grinding is required which may cause changes in the sample. There is no mixing with other reagents as in the preparation of potassium bromide pellets which can introduce water or as in the preparation of Nujol mulls which causes a masking of portions of the spectrum. Further, since no contact with the sample is required sample shape or surface roughness does not

71

Spectroscopy of Rubbers and Rubbery Materials affect the spectra as it does with ATR techniques. The wavelength range covered is limited only by the FTIR instrumentation and the PA detector window used and not by the limitations of ATR crystals. Quantitation is possible even though the theory of photoacoustic spectroscopy is somewhat complex. The main requirement being that the strongest bands not be used or at least checked for saturation. The quality of the results is improved if band ratioing is used since this compensates for instrumental and sample variations. As a surface analytical method the depth being examined can be varied by changing the interferometer mirror velocity. The depth being examined is usually somewhat deeper than that by ATR methods and the techniques are complementary in that respect.

Acknowledgements This work was carried out in support of the research efforts of PPG Industries, Inc., Monroeville Technical Center, Pennsylvania, USA. I wish to thank Denise Callahan and Audrey Anderson of the library staff and Vicki Potter formerly of the secretarial staff for their assistance and Dick Obrycki and Singh Manocha, Heads of the Analytical Departments, and Dave McKeough, Research Director, for their support.

References 1.

A.G. Bell, American Journal of Science, 1880, 20, 305.

2.

A.G. Bell, Philosophical Magazine, 1881, 11, 510.

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R.V. Bruce, Bell: Alexander Graham Bell and the Conquest of Solitude, Little, Brown and Company, Boston, USA, 1973.

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J. Tyndall, Proceedings of the Royal Society of London, 1881, 31, 307.

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W.C. Röntgen, Philosophical Magazine, 1881, 11, 5, 308.

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Lord Rayleigh, Nature, 1881, 23, 274.

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W.H. Preece, Proceedings of the Royal Society of London, 1881, 31, 506.

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M.E. Mercadier, Comptes Rendus Hebdomadaires des Seances de l’Academie des Sciences, 1881, 92, 409.

9.

M.L. Veingerov, Doklady Akademii Nauk SSSR, 1938, 19, 687.

72

Photoacoustic Fouriers Transform Infrared Spectroscopy of Rubbers and Related Materials 10. M.L. Veingerov, Doklady Akademii Nauk SSSR, 1945, 46, 182. 11. A.H. Pfund, Science, 1939, 90, 326. 12. K.F. Luft, Zeitschrift für Technische Physik, 1943, 24, 5, 97. 13. L.B. Kreuzer, N.D. Kenyon and C.K.N. Patel, Science, 1972, 177, 347. 14. E.L. Kerr and J.G. Attwood, Applied Optics, 1968, 7, 5, 915. 15. A. Rosencwaig, Optics Communications, 1973, 7, 4, 305. 16. A. Rosencwaig, Science, 1973, 181, 657. 17. A. Rosencwaig and A. Gersho, Science, 1975, 190, 556. 18. M.M. Farrow, R.K. Burnham and E.M. Eyring, Applied Physics Letters, 1978, 33, 8, 735. 19. M.G. Rockley, Chemical Physics Letters, 1979, 68, 2-3, 455. 20. M.G. Rockley, Applied Spectroscopy, 1980, 34, 4, 405. 21. D.W. Vidrine, Applied Spectroscopy, 1980, 34, 3, 314. 22. J.G. Parker, Applied Optics, 1973, 12, 12, 2974. 23. A. Rosencwaig and A. Gersho, Journal of Applied Physics, 1976, 47, 1, 64. 24. H.S. Bennett and R.A. Forman, Applied Optics, 1976, 15, 10, 2405. 25. L.C. Aamodt, J.C. Murphy and J.G. Parker, Journal of Applied Physics, 1977, 48, 3, 927. 26. F.A. McDonald and G.C. Wetsel Jr., Journal of Applied Physics, 1978, 49, 4, 2313. 27. A. Rosencwaig, Photoacoustics and Photoacoustic Spectroscopy, John Wiley & Sons, New York, NY, USA, 1980. 28. J.M. Chalmers, B.J. Stay, G.F. Kirkbright, D.E.M. Spillane and B.C. Beadle, Analyst, 1981, 106, 1179. 29. K.Y. Wong, Journal of Applied Physics, 1978, 49, 6, 3033.

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Spectroscopy of Rubbers and Rubbery Materials 30. M.J. Adams, A.A. King and G.F. Kirkbright, Analyst, 1976, 101, 1199, 73. 31. M.W. Urban and J.L. Koenig, Applied Spectroscopy, 1985, 39, 6, 1051. 32. M.W. Urban and J.L. Koenig, Applied Spectroscopy, 1986, 40, 4, 513. 33. R.O. Carter III and S.L. Wright, Applied Spectroscopy, 1991, 45, 7, 1101. 34. J.F. McClelland, R.W. Jones, S. Luo and L.M. Seaverson in Practical Sampling Techniques for Infrared Analysis, Ed., P.B. Coleman, CRC Press, Boca Raton, FL, USA, 1993, 107. 35. T. Ando, S. Inoue, M. Ishii, M. Kamo, Y. Sato, O. Yamada and T. Nakano, Journal of the Chemical Society, Faraday Transactions, 1993, 89, 4, 749. 36. C.Q. Yang and W.G. Fateley, Journal of Molecular Structure, 1986, 146, 25. 37. J. Monchalin, L. Bertrand, G. Rousset and F. Lepoutre, Journal of Applied Physics, 1984, 56, 1, 190. 38. S.J. McGovern, B.S.H. Royce, J.B. Benziger, Journal of Applied Physics, 1985, 57, 5, 1710. 39. R.S. Pandurangi and M.S. Seehra, Analytical Chemistry, 1990, 62, 18, 1943. 40. H. Hövel, P. Grosse and W. Theiss, Journal of Non-Crystalline Solids, 1992, 145, 159. 41. N. Teramae, International Polymer Science and Technology, 1985, 12, 6, T/43. 42. N. Teramae, Nippon Gomu Kyokaisi, 1984, 57, 3, 141. 43. R.O. Carter III, M.C. Paputa Peck, M.A. Samus and P.C. Killgoar, Jr., Applied Spectroscopy, 1989, 43, 8, 1350. 44. W.H. Waddell and J.R. Parker, Rubber Chemistry and Technology, 1992, 65, 836. 45. J. McClelland, MTEC Corporation, Ames, IA, USA, Private Communication, 1999. 46. J.P. Coates, Applied Spectroscopy Reviews, 1999, 34, 1-2, 121. 47. R.O. Carter III, Applied Spectroscopy, 1992, 46, 2, 219.

74

Photoacoustic Fouriers Transform Infrared Spectroscopy of Rubbers and Related Materials 48. A. Roychoudhary and P.P. De, Journal of Applied Polymer Science, 1995, 55, 9. 49. G.F. Kirkbright and K.R. Menon, Analytica Chimica Acta, 1982, 136, 373. 50. J.F. McClelland, S. Luo, R.W. Jones and L.M. Seaverson, Proceedings of the SPIE - International Society Optical Engineering, 8th International Conference of Fourier Transform Spectroscopy, Lübeck, Germany, 1991, Volume 1575, 226. 51. J.R. Parker and W.H. Waddell, Journal of Elastomers and Plastics, 1996, 28, 140. 52. J.L. Binder, Analytical Chemistry, 1954, 26, 12, 1877. 53. M. Panyszak and J. Kovar, Canadian Journal of Spectroscopy, 1986, 31, 5, 130. 54. J.M. Sloan and J.P. Clements, US Army Report AMMRC TR 84-25, National Technical Information Service (NTIS), US Department of Commerce, Springfield, VA, USA, 1984. 55. R.S. Silas, J. Yates and V. Thornton, Analytical Chemistry, 1959, 31, 4, 529. 56. ASTM Method D3900-95 (2000) Standard Test Methods for Rubber-Raw Determination of Ethylene Units in EPM (Ethylene-Propylene Copolymers) and EPDM (Ethylene-Propylene-Diene Terpolymers). 57. J.R. Parker, Photoacoustic Analysis of Ethylene-Propylene (EPM) Copolymers, Poster Presentation, ISPAC-10, 10th International Symposium on Polymer Analysis and Characterisation, University of Toronto, Toronto, Canada, 1997. 58. E.V. Thompson in Encyclopedia of Polymer Science and Engineering, 2nd Edition, Ed., J. Kroschwitz, John Wiley & Sons, New York, USA, 1985, Volume 16, 711. 59. I. Furuta, S-I. Kimura and M. Iwama in Polymer Handbook, 4th Edition, Ed., J. Brandrup, E.H. Immergut, E.A. Grulke, A. Abbe and D.R. Bloch, John Wiley & Sons, New York, NY, USA, 1999, Volume 1. 60. M.C. Paputa Peck, M.A. Samus, P.C. Kilgoar, Jr. and R.O. Carter III, Rubber Chemistry and Technology, 1991, 64, 4, 610. 61. P. Delprat and J-L. Gardette, Polymer, 1993, 34, 5, 933. 62. L. Gonon, O.J. Vasseur and J-L. Gardette, Applied Spectroscopy, 1999, 53, 2, 157.

75

Spectroscopy of Rubbers and Rubbery Materials 63. M.G. de Oliveira, O. Pessoa, H. Vargas and F. Galembeck, Journal of Applied Polymer Science, 1988, 35, 7, 1791. 64. K-J. Eichhorn, I. Hopfe, P. Pötschke and P. Schimdt, Journal of Applied Polymer Science, 2000, 75, 1194.

76

3

Infrared Spectroscopy of Rubbers Prajna P. De

3.1 Introduction Infrared spectroscopy is one of the oldest techniques for the molecular level characterisation of materials and it has been extensively used to study polymers. Exactly two hundred years ago Sir William Herschel discovered infrared radiation in 1800. It took another hundred years of research to find a correlation between the rotational and vibrational motion of organic molecules and infrared radiation. Coblentz [1-4] produced a monumental research work, which was published by the Carnegi Institution of Washington in 1905. His work encompassed instrument design, and the development of experimental technique as well as the determination of absorption and reflection spectra of a large number of compounds both organic and inorganic. Infrared (IR) spectrometry was used more in the USA and Europe after World War II, thanks to the availability of high quality spectrometers at reasonable prices, which could produce several spectra per hour and could be operated even by non-experts. The Perkin-Elmer single beam instrument was first on the market and it was soon followed by double beam spectrometer [5], offered by Bechman, Hilger and Gaertner. The instrument was based on a source of infrared radiation, collimating or mirror systems, gratings, detectors and recorders. The absorption or emission of energy in an electromagnetic spectrum occurs in discrete packets of photons. The relationship between the energy of a photon and the frequency appropriate for the description of its propagation is given by the famous Bohr equation: Δ E = hν

(3.1)

where Δ E in ergs represents energy absorbed (or emitted), as a molecule passes from one quantised state to another; ν represents frequency in cycles per second; and h is a universal constant known as Planck’s constant (6.6256 x 10-27 erg-s). For some purposes, it is more convenient to think of radiant energy as a continuous wave motion in which λ represents the interval between the nodes in the wave pattern. The relationship between wave length and frequency is given as:

77

Spectroscopy of Rubbers and Rubbery Materials ν = c/λ

(3.2)

1/λ = ν

(3.3)

and

where λ is the wave length in centimeters and c is the velocity of propagation of radiant energy in a vacuum (speed of light in vacuum is 2.9979 x 1010 cm. s-1 ), ν is the wave number in cm-1. The intensity of a beam of radiation is characterised by its radiant power, which is proportional to the number of photons per second that are propagated in the beam. The various regions in the electromagnetic spectrum are displayed in Table 3.1, along with the nature of changes brought about by radiation [6,7]. The infrared region extends from 2.5 to 50 μ (4000 to 200 cm-1), the region from 12,500 to 4000 cm-1 is called near infrared and the region from 667 to 50 cm-1 is called far infrared. Molecules do not absorb infrared radiation when they have a completely symmetrical charge distribution and there is no change in dipole moment during molecular vibration with different amplitude or rotation at different rates. There are two kinds of fundamental vibrations for molecules:

Table 3.1 Electromagnetic spectral regions useful for chemical analysis Region Name

Approximate Boundaries Wave length

Frequency, Hz

X-ray

10-2 – 10 nm

3x1019 – 3x1016

Vacuum ultraviolet

10 – 200 nm

3x1016 – 1.5x1015

Near ultraviolet

200 – 400 nm

1.5x1015 – 7.5x1014

Visible

400 – 800 nm

7.5x1014 – 3.8x1014

Near Infrared

0.8 – 2.5 μm

3.8x1014 – 1x1014

Fundamental Infrared

2.5 – 50 μm

1x1014 – 6x1012

Far Infrared

50 – 300 μm

6x1012 – 1x1012

Microwave

0.3 – 0.5 m

1x1012 – 6x108

Radiowave

0.5 – 300 m

6x108 – 1x106

78

Infrared Spectroscopy of Rubbers • Stretching, in which the distance between two atoms increases or decreases but atoms remain in the same bond axis, and the angle between a vibrating bond and a chemical bond that is attached to one of the atoms involved in the vibration, is not altered by stretching vibrations. Stretching vibrations are of two types: (i) Symmetrical stretching (ii) Asymmetrical stretching • Bending vibrations are nuclear motions that cause a change in the angle between two vibrating bonds. Bending vibrations are of four types: (i) Scissoring (ii) Rocking (iii) Wagging (iv) Twisting Different kinds of vibrations [8,9] are shown schematically in Figure 3.1.

Figure 3.1 Vibration of a group of atoms

79

Spectroscopy of Rubbers and Rubbery Materials A non-linear molecule that contains n atoms has 3n-6 possible fundamental vibrational modes that can be responsible for the absorption of IR light, whereas a linear molecule has 3n-5 possible vibrations. An approximate value for the stretching frequency (ν in cm-1) of a bond is related to the masses of the two atoms (Mx and My in grams), the velocity of light (c) and the force constant of the bond (k, in N/m): 1

1 ⎡ (M x + M y ) ⎤ 2 ν= ⎢k ⎥ M x M y ⎥⎦ 2πC ⎢⎣

(3.4)

Single, double and triple bonds have force constants that are approximately 500, 1,000 and 1,500 N/m, respectively. Stretching vibrations are found to occur in order of bond strength [9]. The triple bond is stronger than the double bond, which in turn is stronger than the single bond [9]. That is, absorption of the triple bond occurs at 2300-2000 cm-1, while absorption for the double bond occurs at 1900-1500 cm-1 and for the single bond at 1300-800 cm-1. Thus IR spectra are very useful for distinguishing the different functional groups depending upon their mode of vibration. The conception of group frequencies [10,11] is very useful for practical interpretation of IR spectra, and nowadays software is attached to all the modern instruments, providing all the group frequencies.

3.2 Sample Preparation The selection of appropriate sample preparation technique is very important for accurate analysis. Techniques of sample preparation for raw rubber identifications are as follows: i) Preparation of raw rubber film [12] ii) Preparation of film on low density polyethylene (LDPE) matrix [13] iii) Preparation of film prepared from chloroform solution [12] iv) Preparation of solution in chloroform or toluene [12] v) Pyrolysis of raw rubbers [12] Very rarely neat rubbers are used in end-use applications. In general, rubbers are compounded with other materials such as fillers, antioxidants, accelerators and sulfur in 80

Infrared Spectroscopy of Rubbers order to optimise a set of specific properties, e.g., tensile strength, modulus, impact strength, and resistance to oxidation. More often than not, economic factors have an overriding consideration. Carbon black is compounded into a large percentage of all elastomers due to the fact that it is a relatively cheap filler and generally improves the physical properties of rubbers. Hence in order to characterise the compounded and vulcanised rubber products, different nondestructive methods are adopted, other than the pyrolysis of rubber vulcanisates.

3.3 Different Types of IR Spectroscopy Infrared spectra of optically dense materials such as silica and carbon black filled rubbers cannot be obtained in the dispersive infrared spectrometers. In these cases only the destructive methods like pyrolysis [12] of rubber vulcanisates are applied. With the advent of Fourier transform infrared (FTIR) spectroscopy, numerous problems in the field of rubber and polymer characterisation [14-16] could be easily solved due to an increase in the signal to noise ratio, higher energy throughout, data processing capability and rapid scanning where an entire spectrum can be recorded in a matter of seconds. Theoretical considerations of the FTIR spectroscopy are available in the excellent books by Griffiths [17,18] and Ferraro and Basile [19]. Several new sampling techniques allow the polymers to be examined in its fabricated state, i.e., as a powder or fibre (diffuse reflectance), film (reflectance – absorbance), coating (specular reflectance) or bulk sample (photoacoustic). Attenuated total reflectance (ATR) spectroscopy, developed by Harrick [20] and Farenfort [21] is an ideal technique for analysing the surface of crosslinked elastomeric materials [22] and carbon black filled rubber [23]. Diffuse Reflectance Infrared Fourier Transform Spectroscopy (DRIFTS) has recently become a powerful technique for the analysis of powders and coarse solids. The sensitivity of DRIFTS and its quantitative accuracy for powdered samples have been documented by Fuller and Griffiths [24,25]. Similar to ATR spectroscopy, infrared photoacoustic spectroscopy (PAS), originally proposed by Alexander Graham Bell [26] and developed by Rosencwaig and Gersho [27], attached to FTIR has the ability to examine any type of solid, or semisolid material, whether it be crystalline, powder, amorphous, smear, or gel. The surface selective nature of the PAS technique is advantageous for the understanding of surface phenomena. Infrared PAS has been used to study surface treatment of silica [28], and reactions involving silica [29].

3.4 Quantitative Analysis The polymer systems are often complex. For example, crystalline and amorphous regions coexist in semi-crystalline polymers. Any physical or chemical treatment of a polymer will induce structural changes, the knowledge of which is essential for a better 81

Spectroscopy of Rubbers and Rubbery Materials understanding of the polymer properties. Infrared spectra of neat polymer can be taken in solution and quantitative analysis is possible using Beer-Lambert’s law, following the relationship between the concentration of the polymer solution and the amount of radiation absorbed: Log

Io = I

n

∑a c d i i

(3.5)

i =1

where, Io is the intensity of incident radiation, I is the intensity of radiation after passage through the sample, ci is concentration of the component (g/l), d is the thickness of the sample cell in centimetre and ai is the constant called absorptivity or extinction coefficient. Log (Io/I) is called the absorbance. Transmittance (I/Io) of a sample is related to absorbance by the equation: A = log (1/T) = -log T = al ci d

(3.6)

A calibration curve is prepared, using absorbance versus concentration plot, so that the concentration of the unknown component can be determined. But quantitative analysis for a complex system like vulcanised rubber or a blend of two or three components, is not possible. The use of computers with the FTIR spectrometer, increases the rapid scanning capability, data processing for analysis of chemical or physical structural changes in polymers as a function of time over the entire mid-IR frequency. Quantitative analysis of such complex systems by data processing of digitised spectra has been recently developed by Koenig and co-workers [30]. Koenig has developed several methods for quantitative analyses: i) Factor group analysis [31] ii) Ratio method [32,33] iii) Least squares curve fitting [34]

3.5 Applications of IR Spectroscopy Hampton [35] reviewed the application of infrared spectroscopy in the characterisation of rubbers. He discussed the role of infrared spectroscopy in the determination of polymer

82

Infrared Spectroscopy of Rubbers composition, study of molecular structure, tacticity and sequence determination of crystallinity and thermal transitions. Another review [36] discusses the determination of composition and structure, including tacticity, branching and end groups for diene polymers with the help of IR spectroscopy. The infrared spectra for common rubbers have been reported in ASTM D3677 [12]. Most of these spectra were done either in solvent casting film or by pyrolysis. Roy and De [13] used LDPE as the matrix material in the sample preparation, in which the raw rubber is mixed with LDPE at 120 °C in a Brabender Plasticorder and made into a thin film, which is stiff. Peak assignments of some rubbers are given in Tables 3.2a and 3.2b. The peak positions are almost in the same position as reported in ASTM D3677. The representative spectra [13] of natural rubber (NR) and chloroprene rubber (CR) are given in Figures 3.2a and 3.2b. For raw rubber analysis, this method is useful, as it does not need any solvent, and there is no question of destruction of functional groups which happens during pyrolysis. Furthermore, it also characterises the hanging groups of nitrile rubber (NBR) or styrene-butadiene rubber (SBR), unlike the IR-ATR method which deals with only the surface. Composition and microstructure determination of polybutadiene (BR) and natural rubber (NR) can be done by infrared spectra. Three different base units are possible for linear addition polymers of 1,3 butadiene: units with cis or trans internal double bands from 1,4 addition and units with side vinyl groups from 1,2 addition (see Scheme 3.1a). The IR spectra of butadiene polymers have been the subject of considerable study and discussion [37-41]. Early work on the chemical and IR analyses of BR was reviewed and a method described [35,42] for the IR determination of cis, vinyl and trans units from absorption measurements at 724, 911, and 967 cm-1, respectively, as though each sample was a mixture of three isomerically pure cis, vinyl and trans homopolymers (Figure 3.3). Since isomerically pure polymers were not available, three different kinds of BR, each relatively high in one of the three kinds of base units were used as standards [35]. The band near 1308 cm-1 was identified [38,39] with the cis isomer and used for analyses [43]. The 1308 cm-1 band is weak and relatively broad, with the appearance of an unresolved doublet (1306, 1311 cm-1). The cis band at 730 cm-1 is more frequently used in spite of some difficulties. Relatively pure, crystalline stereoregular polymers have been prepared and structures were determined by X-ray diffraction for cis [44], trans [45] and syndiotactic vinyl [46] and isotactic vinyl [47]. Infrared spectra [48-50] have been published for the four stereoregular polybutadienes, with detailed analyses of the spectra and band assignments for cis [51], trans [51] and syndiotactic vinyl [51] polymers. For the spectrum of isotactic vinyl BR, bands at 1232, 1225, 1109, 943, 876, 807 and 695 cm-1 83


Table 3.2a Peak assignments for some commonly used rubbers Name of polymer

Absorption bandsa Absorption of pure polymer bands in LDPE matrix cm-1 films, cm-1

Corresponding chemical group

Assignmentb

CH3 NR

BR

SBR

EPDM (HD)

833 (s)

836(s)

cis C=C

(C-H)op def

1370 (s)

1374 (s)

-CH3

(C-H) def

1665 (m)

1661 (w)

—C = C

(C-C) str

885 (m)

889 (w)

3,4 addition products of isoprene

741 (vs)

745 (s)

cis-CH = CH—

(=C-H)op def

990 (m)

993 (m)

cis-CH=CH-

(=C-H)op def

962 (m)

-

trans -CH=CH-

(=C-H)op def

909 (m)

909 (m)

CH2 = CH-

(CH2)op def

699 (vs), 758 (s)

700 (vs), 760 (s)

Monosubstituted benzene

(C-H)op def

1490 (m)

1493 (m)

C=C in benzene ring

(C-C) str

2000-1660 (m) (4 peaks)

2000-1660 (w) (4 peaks)

diff benzene ring substitution

(C-H)op def

909 (m)

911 (w)

butadiene part

962 (s)

962 (m)

butadiene part

990 (m)

990 (w)

butadiene part

760 (m)

-

-

-

710 (m)

-

-

-

1370 (s)

1365 (m)

-CH3

(C-H) def

1083 (s)

1074 (m)

-CH3-CH2-CH2

-do-

900 (m)

889 (w)

hexadiene

1145

-

C(CH3)2

(C-H) def

a: s, strong; m, medium, w, weak b: str, stretching; tw, twisting; def, deformation; op, out of plane; asym, asymmetric, sym, symmetric. VS: very strong HD: hexadiene Reprinted in modified form from S. Roy and P.P. De, Polymer Testing, 1992, 11, 1, 3, Table 1, with permission of Elsevier Science, 1992, [13]

84

Infrared Spectroscopy of Rubbers

Table 3.2b Peak Assignments for some polar rubbers Name of polymer

CR

NBR

CSM

XNBR

ENR-25

Absorption bandsa of pure polymer films, cm-1

Absorption bands in LDPE matrix cm-1


Assignmentb

1665 (vs)

1660 (s)

>C = C

(C=C) str

1315 (m)

1304 (s)

—CH=CH2

(C-H) def

1110 (s)

1120 (w)

-

-

820 (s)

824 (s)

C-Cl

(C-Cl) str

2220 (vs)

2223 (s)

—C ≡ N

(-C=N) str

1300 (m)

1304 (w)

—CH=CH2

(-C=H) def

1450 (s)

1466 (s)

—CH2

(C-H) def

962 (s)

971 (s)

Butadiene part

909 (m)

909 (s)

-do-

1370 (vs)

1373 (s)

—SO2 Cl

(SO2) asym, str

1265 (m)

1258

—CH2

(CH2), tw

1150 (s)

1160 (s)

—SO2Cl

(SO2) sym, str

740-720

submerged in 723 band on PE

C-Cl

(C-Cl), str

2220 (vs)

2223 (vs)

-C ≡ N

(-C=N) str

1710 (s)

1698 (s)

-COOH

(C-O) str

962 (s)

969 (s)

Butadiene part

909 (m)

916 (w)

-do-

870 (m)

870 (w)

epoxide

Ring vibration

1250 (m)

1244 (m)

epoxide

-do-

835 (s)

836 (m)

C=C(CH3)H

(C=H)op def

1370 (s)

1375 (s)

-CH3

(C-H) def

1665 (m)

1662 (w)

C=C

(C=C), str

85


Table 3.2b Continued Name of polymer

Silicone methyl

Absorption bandsa of pure polymer films, cm-1

Absorption bands in LDPE matrix cm-1


Assignmentb

1373 (w)

1375 (w)

C H2

(C-H) def

1019 (s)

1020 (m)

Si-O

(Si-O) str

1264 (s)

1265 (m)

Si-CH3

(-C-H) sym def

800 (m)

799 (m)

Si-C

(Si-C) str

880 (w)

882 (m)

-CH=CH2

a: s, strong; m, medium, w, weak b: str, stretching; tw, twisting; def, deformation; op, out of plane; asym, asymmetric, sym, symmetric. Reprinted in modified form from S. Roy and P.P. De, Polymer Testing, 1992, 11, 1, 3, Table 2, with permission of Elsevier Science, 1992, [13]

Figure 3.2 (a) Infrared spectra of NR, (b) Infrared spectra of CR Reprinted from S. Roy and P.P. De, Polymer Testing, 1992, 11, 1, 3, Figures 1 and 2, with permission of Elsevier Science, 1992, [13]

86


Scheme 3.1a Structure of polybutadiene base units

Scheme 3.1b Structure of polyisoprene base units

Figure 3.3 Infrared spectra of butadiene polymers Reprinted with permission from R.R. Hampton, Rubber Chemistry and Technology, 1972, 45, 3, 546. Copyright 1972, Rubber Division, American Chemical Society, [35]

87

Spectroscopy of Rubbers and Rubbery Materials have been used to differentiate it from the syndiotactic form. Trans-BR exists in amorphous and two different crystalline forms. Recently Rao and co-workers [52] synthesised 1,3 butadiene using a neodymium chloride-tripentanotate-triethyl aluminium catalyst. FTIR spectrum showed absorption peaks at 740, 912, 965 cm-1 for cis 1,4, vinyl 1,2 and trans 1,4 butadiene but microstructure determination shows that the cis content is about 99%, and thus a stereospecific BR with high cis 1,4 and low vinyl 1,2 can be obtained. Another widely used rubber is NR, which can be present in four base units in linear addition polymers of isoprene (shown in Scheme 3.1b). Infrared analysis of polyisoprene is difficult. The out of plane hydrogen deformations are satisfactory for determining 3,4 units at 888 cm-1 and 1,2 units at 908 cm-1 but the bands at 839 cm-1 for cis 1,4 and 842 cm-1 for trans 1,4 are relatively weak (Figure 3.2a). Richardson and Sacher [53] measured total 1,4 structure at 815 and 857 cm-1, where the broad weak trans and the sharper cis bands show identical absorptivities. Cis and trans proportions were calculated from the absorbance at 843 cm-1, where absorptivity for the cis form is nearly twice as great as for the trans forms. On the basis of X-ray diffraction and chemical evidence, it was thought that NR and balata were 100% cis 1,4 - and 100% trans-1,4-polyisoprene, respectively. Modern infrared and nuclear magnetic resonance (NMR) data strongly support this belief, but there has been an interesting controversy about the possible presence of a small amount of 3,4 isomer in the natural polymers. Binder and Ransaw [54,55] reported the presence of 2.2% 3,4 isomer in NR and 1.3% 3,4 isomer in balata, and no evidence was found for the presence of trans forms in NR, or for cis forms in balata. Following extensive IR studies, Saunders and Smith [56] concluded that gutta percha and balata were trans 1,4 polyisoprene, while Hevea and guayule rubbers were cis 1,4 polyisoprene (spectra shown in Figure 3.2a). Infrared methods for copolymer composition like SBR, NBR, ethylene-propylene copolymer (EPM), ethylene propylene diene terpolymer (EPDM) are listed in Table 3.3. The determination of percentage of styrene and butadiene isomer distribution in copolymers is an extension of the methods for the analysis of polybutadiene. The styrene band at 700 cm-1 is largely independent of the sequence distribution and therefore useful in styrene content determination [76]. A series of bands in the IR spectrum of crystalline isotactic polystyrene at 758, 783, 898, 920, 1053, 1084, 1194, 1261, 1297, 1312 cm-1 have been attributed to the helical structure [77]. The absorption bands for butadiene in SBR are similar to BR structures (Table 3.2a). Like styrene content determination, acrylonitrile content in nitrile rubber can be determined with IR spectra following the quantitative analysis method. Butadiene units, adding up to an acrylonitrile unit in the polymer chain, show a strong tendency to add 1,4 trans structure. The acrylonitrile NBR has the following peaks:

88

Infrared Spectroscopy of Rubbers 2920, 2851 cm-1 for ν (-C-H of CH3-), CH22222 cm-1 for ν (-C≡N) 1460 cm-1 for δ(C-H) of -CH2-, 1304 cm-1 for γ(CH2) of -CH2 962 cm-1 for δ(=C-H)op of trans -CH=CH909 cm-1 for δ(CH2)op of CH2 = CH740 cm-1 for δ(=C-H) op of cis - CH = CH- (Table 3.2b).

and

Table 3.3 Example of quantitative infrared analysis from copolymers Determination of polymer/rubber

Sample analysed

Infrared Technique

Reference

BR, blends

AT R

57

Styrene

SBR

Film

58

Styrene

SBR

Film

59

Acrylonitrile, styrene

Copolymer

Solution

60

Acrylonitrile, styrene

Copolymer

Film

61

Butadiene, methyl methacrylate

Copolymer

Solution

62

SBR-NR vulcnaisate

Pyrolysis

63

Blends

Pyrolysis

64

Oil blends

Pyrolysis

65

SBR

NR-SBR blend

Pyrolysis

66

NR, NBR, SBR

Binary blends

Pyrolysis

67

IIR, Polypropylene

Blends

Pyrolysis

68

NR, SBR, BR

Blends

Pyrolysis

69

Tyre tread

Film

70

Propylene

EPM

Solution

71, 72

Propylene

EPM

Film

73

EPDM

Film

74, 75

Cis, trans, vinyl

NR NR, SBR, IIR NR, SBR

BR

Dicyclopentadiene

Reprinted in modified from [35] with permission of Rubber Chemistry Division, American Chemical Society, USA, 1972

89

Spectroscopy of Rubbers and Rubbery Materials An IR spectrum shows the butadiene units to be almost exclusively trans [78], (that is, 98% trans and 2% vinyl, and 0% cis). Propylene content in EPM rubber can be determined with the help of IR spectra. A propylene band near 1155 cm-1 has been widely used [79] for EPM analysis, frequently in combination with the polyethylene band at 721 cm-1. Tacticity is important in EPM rubber, and the bands at 1229 and 1252 cm-1 are characteristic of syndiotactic and isotactic structures, respectively, (both bands are present in atactic polypropylene as well). Polymer structure may vary in the relative tactic placement of adjacent head to tail propylene units and in the sequence distribution of base units along the chain. Some of them can be identified [80] by infrared spectra, such as isolated or head to tail propylene units: —CH2 –CH (CH3) –CH2-, at 1155 cm-1; head to head propylene units, —CH2 –CH (CH3)–CH (CH3) –CH2 at 1120 cm-1; and side ethyl groups, –CH—(CH2 - CH3)- at 775 cm-1 can be used to identify crystalline ethylene blocks. Depending on the conditions of polymerisation the distribution of monomer units in ethylene – propylene copolymers may be random or may show a preference for the formation of relatively long sequences (blocks) of the same base unit [81], or may show a tendency for alternation [82] of ethylene and propylene units. The distribution of base units can be calculated theoretically [83] from monomer reactivity ratios and can be measured experimentally by IR spectroscopy. For random copolymers with propylene content, most of the propylene units were found in head to tail sequences and thought to be isotactic [84]. An EPDM rubber is produced by the terpolymerisation of ethylene and propylene with a small amount (typically of the order of 5%) of an unconjugated di-olefin. The di-olefins used, include dicyclopentadiene, 1,4-hexadiene, 5-methylene-2-norbornene, 5-ethylidene2-norbornene and methyl tetrahydroindene, 1,5 cyclo octadiene. A number of other dienes [74,75] have been tried. Infrared spectroscopy [35] is used to find out the ter monomer content. The characteristic peaks for the ter monomer are shown in Table 3.4. In view of the relatively low concentrations, it is probable that ter monomer base units are present largely as isolated units in EPDM but the distribution of propylene and methylene sequences is of considerable interest. CR is a polymer with relatively large monomer repeating units that consequently exhibit little vibrational coupling between the chemical units along the chain [85]. Infrared analysis [86] showed the following bands: cis 1652 cm-1, trans 1660 cm-1; 1,2, 925 cm-1; 3,4, 883 cm-1. For emulsion polymerisation addition is almost entirely 1,4 with no more than about 2% 1,2- and 3,4-. Cis 1,4 – CR exhibits characteristic infrared absorption bands at 847, 1652, 3025 and 3282 cm-1 (C = C overtone). The corresponding trans bands are at 822, 1660, 3018 and 3295 cm-1. Typical CR has 78%-96% trans, 1,4; 18% cis 1,4,

90


Table 3.4 Characteristic olefin infrared frequencies for EPDM Hydrogen stretch cm-1

Double bond cm-1

Hydrogen stretch cm-1

3043

1611

600

-

-

965

5-methyl-2-norbornene

3069

1662

872

5-ethylidene-2-norbornene

3040

1688

809

Termonomer Dicyclopentadiene 1,4 hexadiene

Reprinted from [35] with permission of Rubber Chemistry Division, American Chemical Society, USA, 1972

0.3-2% 1,2-; and 0.2-2% 3,4- compositions. Infrared band at 780 and 950 cm-1 are associated with the crystallinity, normally present in chloroprene rubber at room temperature [87]. Coleman and co-workers [88] determined the proportion of the crystalline phase of CR using the subtraction method. This technique considerably simplified the assignment of the individual IR bands arising from the crystalline and amorphous components of the spectrum [88]. The crystallinity in CR is affected by the temperature of polymerisation. Polymers prepared in the range of –20 °C to 40 °C revealed the specific crystalline bands, which were sensitive to a number of structural irregularities. It is well known that as the polymerisation temperature is increased, the concentration of the following units increased [89]: head to head, trans 1,4; cis 1,4; 1,2; 3,4 units. Coleman and Painter [90] had made detailed studies of structural changes, crystallinity changes of CR with temperature. The infrared spectrum [13] of CR is given in Figure 3.2b showing the absorption bands at 1660, 1304, 1120, 824, 720 cm-1, characteristics of trans isomers (Table 3.2b). Polyurethanes (PU) are a special class of polymers containing the urethane group, (—NH-CO-O-), commonly prepared by a condensation reaction between a diisocyanate and a diol, such as hydroxyl terminated polyether or polyester: HO –R —OH + OCN –R′ –NCO ( O —R —O —CO –NH –R′ –NH –CO ) n~ The isocyanate (2270 cm-1) uretedinedione ring carbonyl (1780 cm-1) and urea carbonyl (1660 cm-1) groups can usually be identified. Carbonyls from ester, urethane, allophanate, isocyanuric acid ring and Biuret groups all absorb near 1730 cm-1 and are difficult to distinguish. Hydrogen bonds which can function as physical crosslinks in PU have been

91

Spectroscopy of Rubbers and Rubbery Materials studied for model compounds and polymers [89,91]. The 3305 cm-1 NH band and inflections at 1780 and 1724 cm-1 on the 1735 cm-1 urethane carbonyl band are attributed to hydrogen bonding; the free NH band is at 3340 cm-1. The NH group may bond to the urethane carbonyl or to the ester carbonyl of polyester urethanes or to the ether oxygen of polyether urethanes [89].

3.5.1 Rubber Blends Polymer blends have had considerable attention recently due to the ease with which the polymer properties can be manipulated to achieve some special characteristics that cannot be achieved by the single polymer systems, or by simple copolymerisation. Different polymers are blended to prepare the various types of industrial products, ranging from giant tyres to small seals. Polymer blends may be immiscible, miscible, compatibilised, reactive or self crosslinking. Utracki [92] described the criteria for the formation of various types of blends. To study polymer miscibility, IR spectroscopy has emerged as one of the most useful analytical tools. The technique can elucidate specific interactions, such as hydrogen bonding, dipolar interaction, etc., [93,94]. The use of IR spectra for characterisation of the blends of common rubbers (NR, BR, SBR) are shown in Table 3.3. The production of reactive polymer alloys and blends has become a frontier area of research and developmental activity in polymer technology in the past two decades. Chemical modification of the polymer backbone, grafting onto a polymer chain, interchain reactions and formation of interpenetrating networks (IPN) are the subject of a number of reviews [95,96]. Maleation [97,98] is one of the oldest and most widely used techniques for modification of polymers. It is the grafting of maleic anhydride onto a polymer backbone by means of radical initiators such as peroxide. Maleation also facilitates specific interactions that improve polymer-polymer compatibility [99]. In order to understand the mechanism of grafting, dibutyl maleate was grafted on to polyethylene (PE) and EPDM using 0.5 wt% to 0.2 wt% dicumyl peroxide (DCP) in a Brabender Plasticorder in two steps at 110 ± 2 °C and 180 °C [100]. The grafted polymers (Peg dibutyl maleate (DBM) and EPDMgDBM) were characterised by IR analysis. The IR spectra [100] of PE, PEgDBM and EPDMgDBM are presented in Figure 3.4. In the spectra of functionalised polymers, the two bands at 1738 cm-1 and 1164 cm-1 were attributed to C=O stretching and C—O stretching of the ester group, respectively. The carboxyl band at 1738 cm-1 was used to determine the degree of grafting. It is of considerable technological importance to make heat, oil and fire resistant (HOFR) cable sheathing compounds by combining the flame and oil resistance of polychloroprene (PCP) and heat and ageing resistance of EPDM through application of judicious techniques of blending. But PCP is not compatible with EPDM. Sen and co-workers [101] prepared

92


Figure 3.4 Infrared spectra of (a) PE, (b) PEgDBM, (c) EPDM, (d) EPDMgDBM Reprinted with permission from A.K. Sen, B. Mukherjee, A.S. Bhattacharyya, P.P. De and A.K. Bhowmick, Die Angewandte Makromolekulare Chemistry, 1991, 191, 3206, 15. Copyright 1991, Wiley-VCH [100] with permission of Huthig and Wepf Verlag Basel, 1991

low halogen fire resistant low smoke (FRLS) cable sheathing compounds from the blends of polyvinyl chloride (PVC) and polyolefins using dibutyl maleate grafted polyolefins as compatibiliser and characterised their properties. PCP/EPDM blends were prepared using dibutyl maleate grafted ethylene-propylene-diene rubber [102]. The probable mechanism of compatibilisation [102] between PCP and EPDMgDBM is shown in Scheme 3.2. EPDMgDBM gives rise to an interaction, mainly through a hydrogen bond between the carboxyl group of DBM and methine hydrogen of PCP. Besides this, dipole-dipole interactions of type >C=O…..Cl–C may also occur [102]. Peroxide initiated graft copolymerisation of vinyl trimethoxy silane (VTMO) and vinyl triethoxy silane (VTEO) onto PE and ethylene propylene copolymer (EPR) was also studied by De and co-workers [103]. The absorbance of —Si –O –Si – group stretching of the grafted silane at 1090 cm1 was used for characterisation of PEgVTMO and PEgVTEO. Actually, in the presence of moisture and a condensation catalyst, the alkoxy group of the silane derivatives pendent to the grafted co-polymer is converted into silanol groups which undergo a condensation reaction to form a –Si – O – Si – type network [103]. Both the hydrolysis of alkoxy silane to silanol and their condensation reaction occur almost instantaneously forming the

93


Scheme 3.2 Probable mechanism of compatibilisation between EPDMgDBM and PCP. Reprinted with permission from A.K. Kalidaha, A.S. Bhattacharya, A.K. Sen and P.P. De, Die Angewandte Makromolekulare Chemistry, 1993, 204, 3484, 19. Copyright 1993, Wiley-VCH, [102]

network [103]. These types of reactive blends are generally used for making special purpose products for example butyl rubber (IIR) is blended with polyamide 6,66,610 copolymer (LPA) in the presence of alkyl phenol formaldehyde resin (PR) to produce IIR-PR-LPA co-polymer [104] which has a novel gas-barrier property, used for rubber hose in an air conditioning system. From Figure 3.5, it is clear that following the reaction of LPA with PR, the peaks at 1548 cm-1 (NH bending) and 1647 cm-1 (C=O) stretching of LPA shifted by the reaction to 1542 and 1640 cm-1, respectively. The peak intensity of the NH bending deformation of the amide group at 1548 and 3085 cm-1 relative to that of CH2 stretching at 2950 cm-1 decreased and a peak of CH2OH of the PR-LPA copolymer at 3372 cm decreased a little from PR. This indicates that the hydroxyl group in PR reacts with the amide group in LPA to produce the co-polymer. IR spectra of the IIR-PR-LPA copolymer and IIR are shown in Figures 3.5e, and 3.5d. In Figure 3.5e, the peak intensity of the double bond at 890 cm-1 and that of CH2OH at 3372 cm-1 decreased, and a new peak at 1750 cm-1 appears, which may be due to a Chroman-ring type of structure indicating that a Diels-Alder type of reaction takes place between PR-LPA copolymer with IIR.

94


Figure 3.5 IR spectra of (a) polyamide 6,66,610 (LPA), (b) alkyl phenol formaldehyde resin (PR), (c) reaction product PR/LPA, (d) IIR, (e) IIR-PR-LPA copolymer Reprinted with permission from S.I. Goto, K. Kimura, T. Yamamoto and S. Yamashita, Journal of Applied Polymer Science, 1999, 74, 3548. Copyright 1999, John Wiley & Sons, Inc., [104]

Another such example is EPDMgNVP (ENVP), a new graft co-polymer which was synthesised to improve the dyeability of EPDM [105]. ENVP was synthesised by grafting N-vinyl pyrrolidione (NVP) with benzoyl peroxide as an initiator in toluene. The IR spectra of ENVP exhibited a characteristic absorption band at 1670 cm-1 (C=O stretching) due to NVP and 1460 cm-1 (-CH2 bending) and 1375 cm-1 (-CH3 bending) due to EPDM. From the literature, many examples of such grafted and compatibilised blends can be found where IR spectroscopy has been used to characterise the blends. Examples of such blends are as follows:

95

Spectroscopy of Rubbers and Rubbery Materials carboxyterminated polybutadiene-nylon [106], polydimethyl siloxane (PDMS) – ethylene methyl acrylate [107,108], epoxidised NR (ENR) – PE co-acrylic acid [109], PU-polyimide [110], ethylene vinyl acetate – maleic anhydride [111], acrylate rubber – polyamide [112], polybutadiene – CR [113], PU – aminoethyl aminopropyl PDMS [114], chlorohydrin rubber grafted with polybutyl acrylate [115].

3.5.2 Self-crosslinking Blends It has been already mentioned that functionalisation of polymers results in new materials with a wide spectrum of properties not available in the parent polymers. Chemical modification of the polymer backbone, grafting onto a polymer chain, interchain reactions and formation of IPN are the subject of a number of reviews [116]. De and co-workers studied a series of self-crosslinkable polymer blends, which are mixtures of two or more functionally reactive polymers that are capable of undergoing mutual crosslinking via condensation or substitution reactions at high temperatures. Such self-crosslinkable rubber-rubber blends include ENR and carboxylated nitrile rubber (XNBR) [117,118]. ENR was found to form self-crosslinkable blends with many chlorine containing polymers like CR [119], PVC [120] and chlorosulfonated polyethylene (CSM) [121-124]. Mukhopadhyay and co-workers [124] reported that CSM reacts with XNBR forming a self vulcanisable miscible blend through the formation of a sulfonyl ester group (Scheme 3.3a). Roychoudhury and co-workers [125] made a detailed study of the reaction with FTIR-ATR spectra and showed that crosslinking takes place in two ways: either through allyl chloride [Scheme 3.3b (1,2)] or through allyl chloride and amide [Scheme 3.3b (3,4)] produced by hydrolysis of a –C≡N group. The IR spectra of a self-crosslinked blend of CSM/XNBR [125] following the reaction scheme (Scheme 3.3b) are presented in Figure 3.6. It is evident that absorption at 1740 cm-1 in the crosslinked blend (Figure 3.6b) is likely to correspond to the C=O stretching of ester. The band at 1124 and 1171 cm-1 may be ascribed to C-O-C asymmetric stretching vibration of an aliphatic ether, but formation of the ester is much stronger, because the band in the region of 1100-1200 cm-1 can also be attributed to C-O-C stretching vibration of an aliphatic ester. The bands at 1697, 968, 916 cm-1 are due to hydrogen bonded –COOH, trans –CH=CH- and CH2 = CH~ (out of plane bending). The absorption band at 2237 cm-1, which is attributed to a C≡N stretching mode is strongly attenuated by self-crosslinking of the blend. This indicates a significant loss or conversion of cyanide groups, due to cyanide/cyanide [126] or cyanide/ carboxyl [127] reactions or even hydrolysis of the cyanide group (shown in Scheme 3.3b). The band at 1553 cm-1 is due to an amide II band and the band at 1651 cm-1 is due to

96


Scheme 3.3a Reaction between Hypalon and XNBR Reprinted with permission from S. Mukhopadhyay, P.P. De and S.K. De, Journal of Applied Polymer Science, 1991, 43, 347, Figure 3. Copyright 1991, John Wiley and Sons, Inc., [124]

amide I, while the band at 968 cm-1 (trans –CH=CH-) is attenuated in the crosslinked blend, and the band at 725 cm-1 is a composite band of cis 1,4 –CH=CH- and CH2 rocking at 740-720 cm-1. Hence cis-trans isomerisation is also exhibited during the crosslinking of blend. Other examples of such self-crosslinkable blends are: ENR/XNBR [128], ENR/CR [119], ENR/XNBR/CR [128,129], ENR/CSM [130], NBR/PVC [131], HNBR/PVC [132], ENR/PVC [133], PVC/ENR/XNBR [134], PAA/ENR [135], XNBR/ CIIR [136]. The blends have been characterised by IR spectroscopy.

97


Scheme 3.3b Reaction between CSM and XNBR Reprinted with permission from A. Roychoudhury and P.P. De, Journal of Applied Polymer Science, 1997, 63, 1761. Copyright 1997, John Wiley and Sons, Inc.

3.5.3 Polyurethanes The properties of PU depend on the proportions of hard and soft segments in the polymer structure. Paul and co-workers synthesised segmented block copolymers of NR and 1,3 butanediol – toluene diisocyanate oligomers [137], NR and bisphenolA-toluene diisocyanate oligomers [138] and characterised them by IR spectroscopy. The

98


Figure 3.6 FTIR-ATR spectra of (a) uncrosslinked 50/50 (w/w) CSM/XNBR blend, (b) self-crosslinked 50/50 (w/w) CSM/XNBR blend, (c) (b-a) difference spectrum Reprinted from A. Roychoudhury and P.P. De, Journal of Applied Polymer Science, 1997, 63, 1761. Copyright 1997, John Wiley and Sons, Inc., [125]

99

Spectroscopy of Rubbers and Rubbery Materials stoichiometric amounts of hydroxyterminated liquid natural rubber (HTNR), 1,3 butanediol (BDO), dibutyl tin dilaurate (DBTDL) and toluene diisocyanate (TDI) in tetrahydrofuran (THF), were mixed following a one shot or two shot process to prepare the block copolymer, which were analysed by IR spectroscopy. Hydrogen bonding plays a critical role in determining the morphology and overall properties of PU, PU-ureas, and other polymers which have pendent functional groups capable of forming hydrogen bonding. The extent of hydrogen bonding in PU can be qualitatively studied by determining the frequency shifts in hydrogen bonded (-N-H) and (-C=O) peaks (that is, -N-H…O=C) relative to the free (-N-H) and (C=O) peaks. The infrared absorption bands for such hydrogen bondings are 1703-1710 (C=O….. H-N), 1730-1740 (free C=O, urethane) 1660-1670 (C=O….. H-N, urea disordered), 1630-1645 (C=O……H-N, urea ordered), such peaks are observed in PDMS and polyether based urethane-urea copolymers [139], PU based on hydroquinone – bis (β-hydroxyethy)ether [140], propylene oxide/ polybutadiene bi-soft segment urethane/urea prepared from polypropylene oxide based ioscyanate – terminated triol prepolymer with polybutadiene diol [141]. The FTIR studies were conducted on UV curable films of water borne polyurethane-acrylate (PUA) ionomer [142] and cyclo aliphatic di epoxide-ENR-glycidyl methacrylate (GMA) [143]. The IR spectra of a cycloaliphatic epoxide-ENR-GMA system based coating (Figure 3.7a) [143] exhibited the following characteristic absorption bands: • Epoxy ring stretching at 1255 cm-1, • Terminal oxiranes of GMA at 909 cm-1, • Weak symmetry in plane ring deformation of internal oxirane at 814-815 cm-1, • Cycloepoxy stretching at 789 cm-1, 1638 cm-1, 844-858 cm-1 due isoprene double bonds of ENR-50, • Acrylic double bonds of GMA at 1638 cm-1 and 815 cm-1. A marked increase in OH stretching at 3550 cm-1 on UV exposure (Figure 3.7b) with concurrent decrease in the absorption bands corresponding to 1255 cm-1 (epoxy, ring stretching) 909 cm-1, 814-815 cm-1 and 789 cm-1 establishes the successful accomplishment of ring opening polymerisation. The increase in 1089 cm-1 can be attributed to the formation of an ether bond. The marked decrease in the absorption bands of isoprene double bonds at 1638 and 844-858 cm-1 (isoprene C-H wagging) showed that free radicals generated during the photolysis of the initiator cause simultaneous radical polymerisation involving the isoprene double bonds of ENR and the acrylic double bonds of GMA. The absorption band at 909 cm-1, corresponding to terminal oxiranes of GMA almost

100


Figure 3.7a FTIR spectrum of cycloaliphatic epoxide-ENR-GMA hybrid system

Figure 3.7b FTIR spectrum of cycloaliphatic epoxide-ENR-GMA hybrid system after UV exposure and post cure Reprinted with permission from R.N. Kumar, C.K. Woo and A. Abusamah, Journal of Applied Polymer Science, 1999, 73, 1569. Copyright 1999, John Wiley and Sons Inc., [143]

101

Spectroscopy of Rubbers and Rubbery Materials disappears on UV irradiation (Figure 3.7b), suggesting that the epoxy ring participates in polymerisation reactions [143]. The results showed that the acrylic double bonds as well as the epoxy groups of GMA, the isoprene double bonds and epoxy groups of ENR and epoxy groups of cycloaliphatic epoxide resin all participate in photoinitiated polymerisation and cross linking reactions, producing an interpenetrating network (IPN).

3.5.4 Rubber-filler Interaction The reinforcement of an elastomer by a filler is associated with a strong interaction between the filler surface and the elastomer [144]. Although a complete knowledge of the exact nature of elastomer/filler interaction is still lacking, the polymer/filler attachments appear to be both physical and chemical, depending on the physicochemical character of the filler surface and the chemical nature of the elastomer [145]. IR spectroscopy in its various forms is an important and forceful technique, which can provide useful information about surface functional groups of carbon black and silica filler. The first IR transmission spectrum of a carbon black with a high oxygen content was reported by Hallum and Drushel [146] using a Nujol mull technique. Donnet [147] and Papirer [148] examined the surface groups on a series of carbon blacks, using direct transmission and internal reflection spectroscopy and observed the presence of bands at 3200-3400, 1740, 1600, 1275-1195 cm-1 indicating hydroxyl, carboxyl, lactone, quinone, ether, phenol groups. Koenig and co-workers [149] utilised IR spectra to find the reversion resistance in NR curing when filled with different grades of carbon black. They found the transmethine (965 cm-1) content is higher at low cure times but less at longer times for the carbon black filled samples. Roychoudhury [150] while studying the interaction of ENR with carbon blacks (N121 and oxidised N121), used diffused reflectance IR spectra to study the surface groups present on the carbon black surface. The spectrum corresponding to the oxidised sample was characterised by more intense peaks in the O-H and the C=O stretching regions, which were responsible for chemical interaction between ENR and oxidised carbon black. De and co-workers [151] using FTIR spectra characterised the bound rubber formation between CSM and non-oxidised and oxidised N110 carbon black. PAS-FTIR spectra [151] (Figure 3.8) of pure CSM, the carbon black (N 110) and CSM/N110 were taken under high vacuum. The spectra showed that carbon black filled CSM compound retained almost all the characteristic bands (1736, 1697, 1651, 1558, 1458 cm-1) due to carbon black, but the strong absorption bands at 1161, and 1366 cm-1, characteristic of –SO2Cl groups of CSM diminished (Figure 3.8c) forming a coupling bond between the CSM macroradicals and quinone or phenoxy radical sites on oxidised carbon blacks, giving an ether type of bond (Figure 3.8c), e.g., CB—)—O-CH~ (1061,1096 cm-1), according to the following reaction:

102


Figure 3.8 FTIR photoacoustic spectra of (a) net CSM, (b) carbon black N110, (c) CSM/N110 compound Reprinted with permission from A. Roychoudhury and P.P. De, Journal of Applied Polymer Science, 1995, 55, 9. Copyright 1995, John Wiley and Sons, Inc., [151]

103

Spectroscopy of Rubbers and Rubbery Materials CB —) OH + ~ CH~ | SO2Cl→CB. —) –O—CH ~+ SO2 + HCl

(3.7)

Hence surface oxidation of carbon black leads to a greater extent of bound rubber formation in polar rubbers [152] like ENR. Chemical interaction between furnace blacks and rubbers is responsible for the greater extent of reinforcement for polar rubbers [153]. Lately, it has been reported that –OH and –COOH groups of surface oxidised furnace blacks react chemically with functionalised polymers like CSM [154], XNBR [155] and ENR [156], forming a crosslinked vulcanisate in the absence of any crosslinking agent. The mechanism of chemical interactions is supported by IR spectroscopy. It is a well known fact that coupling agents cause improvement in rubber-filler interaction in the case of mineral fillers. Recently Bandyopadhyay and co-workers [157] showed that 3-amino propyl triethoxy silane acts as a promoter in bonding between XNBR and surface oxidised carbon black. Manna and co-workers [158] made detailed studies on the effects on the surface oxidation of filler and silane coupling agent, namely N-(4-vinyl benzyl) —N′—[3-(trimethoxysilyl)propyl] ethane-1,2-diamine mono HCl salt (Trade name Z6032; made by Dow Corning, USA) on the properties of ENR filled with intermediate super abrasion furnace (ISAF) carbon black in the absence of conventional crosslinking agents. The various techniques of IR spectra were used to find out the chemical interaction between ENR and oxidised ISAF black [158] and they are shown in Figures 3.9a and 3.9b. The DRIFTS of carbon black sample [158] (Figure 3.9a) show peaks at 1350, 1385 cm-1 due to symmetric and asymmetric C-O stretching of phenol [147], 1574, 1634 cm-1 for tetrahydroanthroquinone and polycyclic quinone [148], 1750, 1800 cm-1 for unsaturated lactones [148], 3400, 3640, 3760 cm-1 for free OH groups. The main difference between two grades of carbon black is the presence of carboxylic acid groups at 1717 cm-1 for oxidised carbon black. The ATR-FTIR spectra of ENR and transmission spectra [158] of thin film of the coupling agent are shown in Figure 3.9b, in which the band at 1085 cm-1, is characteristic of Si—OCH3 stretching vibration, the band at 1602, 1627 cm-1 are characteristic of ~NH2+ deformation modes and C=C (vinyl benzene) stretching, respectively. The peak at 820 cm-1 is due to the rocking vibration of ~NH2+~, whereas the peaks at 1193, 715, 450 cm-1 are characteristic of secondary aliphatic amines, and the peaks at 878, 840 cm-1 are characteristic of epoxy ring vibration, while the peaks at 1050, 1114 cm-1 are

104


Figure 3.9a DRIFTS of carbon black (a) ISAF black, (b) oxidised-ISAF black Reprinted with permission from A.K. Manna, P.P. De, D.K. Tripathy, S.K. De and M.K. Chatterjee, Rubber Chemistry and Technology, 1999, 72, 2, 398, Figure 1. Copyright 1997, Rubber Division, American Chemical Society, Akron, [158]

due to C=O stretching of ENR, and at 1377 cm-1 and 1648 cm-1 are due to C-CH3 and C=C stretching of ENR. The IR spectra of ENR mixed with ISAF black, oxidised ISAF black, with and without coupling agent are presented in [158], and representative spectra are given in Figure 3.10, in which the spectra of ENR with oxidised ISAF black (60 phr) and coupling agent, Trade Name Z6032 (4 phr) with different moulding times show that the intensity of the peak 874 and 840 cm-1 due to epoxy ring vibration decrease.

105


Figure 3.9b Transmission FTIR spectra of (a) coupling agent, (b) ENR Reprinted with permission from A.K. Manna, P.P. De, D.K. Tripathy, S.K. De and M.K. Chatterjee, Rubber Chemistry and Technology, 1999, 72, 2, 398, Figure 2. Copyright 1997, Rubber Division, American Chemical Society, Akron, [158]

The difference spectrum shows negative absorbance at 1720 cm-1 which is due to the reaction between –COOH groups on the oxidised carbon black surface and to –NH2+— and Si-OCH3 of silane coupling agent. This is further substantiated from the appearance of a peak at 1635 cm -1 with a shoulder at 1620 cm -1, which is due to formation of –(C=O)—N—(amide) and –Si-O-(C=O)- (silyl ester). The surface of hydrated (precipitated) silica is highly polar and hydrophilic due to its numerous silanol groups. The polar surface chemistry assists the development of better

106


Figure 3.10 (a) ATR-FTIR spectra of mix EOZ4 (after 1 h moulding) (b) ATR-FTIR spectra of mix EOZ4 (after 1 min moulding) (c) Difference spectrum (a-b) Reprinted with permission from A.K. Manna, P.P. De, D.K. Tripathy, S.K. De and M.K. Chatterjee, Rubber Chemistry and Technology, 1999, 72, 2, 398, Figure 4. Copyright 1999, the Rubber Division, American Chemical Society, Akron, [158]

107

Spectroscopy of Rubbers and Rubbery Materials bonding of the silicas with polar polymers such as NBR [159], XNBR [160] and CSM [161]. Koenig and co-workers [162] studied the interactions of silica filler in zinc activated sulfur vulcanised cis-1,4 polyisoprene using FTIR-ATR spectroscopy. They observed that an increase in the silica level resulted in band broadening and a frequency shift to lower wave numbers of the silica absorbance region (1250-1000 cm-1) due to a combination of physical and chemical adsorption of the rubber and parts of the cure system on the silica surface. The peaks that appeared near 1040 and 1017 cm-1 were attributed to physical interactions between silica and NR. Manna and co-workers [163] studied the role of the coupling agent Z6032 in bonding between precipitated silica and ENR during high temperature moulding at 180 °C. The ATR-FTIR spectra of EHZ5 (ENR 100, silica 60, coupling agent 5 phr) at different moulding times, are shown in Figure 3.11. From the difference spectrum, it is evident that the peaks at 876 and 3700 cm-1 show negative absorbance. The results signify that a portion of the epoxy groups react during moulding (Figure 3.11c). The new peak at 485 cm-1 is characteristic of C-N-C (tertiary amine) deformation, whereas the positive absorbance at 1078, 1011, 796 cm-1 are characteristic of the Si-O-Si (asymmetric stretch), -Si(OCH3)2 deformation, indicating that the epoxy group reacts with amine functionality to form C-N bonds. The negative peak at 1647 cm-1 confirms that C=C bond reacts with hydrogen chloride (HCl) liberated from the coupling agent during high temperature moulding. The positive absorbance at 704 cm-1 and 1260 cm-1 confirms the formation of C-Cl bonds, and secondary alcohols. In presence of coupling agent, coupling bonds [163] are formed through functional groups (amine and methoxy) in addition to silyl ether linkages as shown in Scheme 3.4. Carbon-silica dual phase fillers (CSDPF), a new generation of materials for rubber reinforcement, have been analysed by Murphy and co-workers [164] with IR spectra (potassium bromide plate method). The region above 3000 cm-1 contains information about the silanols, but unfortunately for carbon black and the dual phase fillers, this information is lost due to interference from a large water band. The two major bands are located, one at about 1600 to 1625 cm-1 and the other at 1115 cm-1. The 1600-1625 cm-1 has been assigned to bulk carbon and adsorbed water, when the 1100 cm-1 region has been ascribed to SiO2 networks, a small peak near 1200 to 1250 cm-1 is indicative of carbon phase oxygen groups.

3.5.5 Milling The milling of rubber is an important operation for further processing of rubbers. During milling, rubber molecules break down, and the chemical changes which occur, are a function of time, temperature, rate of shear, shear stress and viscosity of rubber. The mill

108


Figure 3.11 ATR-FTIR spectra of mix EHZ5. (a) Sample moulded for 60 minutes (b) Sample moulded for 2 minutes (c) The difference spectrum (a-b) Reprinted with permission from A.K. Manna, P.P. De, D.K. Tripathy, S.K. De and D.G. Peiffer, Journal of Applied Polymer Science, 1999, 74, 389, Figure 7. Copyright 1999, John Wiley and Sons, Inc., [163]

breakdown behaviour of NR [165,166] and ENR [167] have been reported earlier, but Kumar and co-workers [168] used IR spectroscopy to detect the formation of functional groups during the milling of NR, ENR, SBR, CR and acrylic rubber. The concentration of carboxyl (>C=O) groups at 1700-1800 cm-1 increases with time of milling for all the

109


Scheme 3.4 Probable mechanism of bonding between ENR and precipitated silica in the presence of silane coupling agent Reprinted with permission from A.K. Manna, P.P. De, D.K. Tripathy, S.K. De and D.G. Peiffer, Journal of Applied Polymer Science, 1999, 74, 389, Figure 9. Copyright 1999, John Wiley and Sons, Inc., [163]

110

Infrared Spectroscopy of Rubbers rubbers. The concentration of hydroxyl groups (—OH) in NR and ENR also increases during milling, but the concentration of epoxide rings decreases and that of furan rings increases with time of milling.

3.5.6 Adhesion IR spectroscopy plays a great role in the study of the chemical changes which occur during the adhesion of polymers with metal. For example, FTIR has been used to analyse the interface of laminates [169] composed of ethylene copolymers and aluminium. The spectra of the thin solution cast films of carboxylic acid copolymer showed two bands in the carbonyl region that were assumed to be due to free and dimeric acid groups, with a free acid group at 1740 cm-1. The increased amount of free acid at interface can be attributed to the interfacial hydrogen bonds between the hydroxyl groups of the acid and the surface AlOH groups or eventually Al-O-Al, thus setting some of the carbonyl groups free [169]. Tack is the ability of two rubbery materials to resist separation after bringing their surfaces into contact for a short time under light pressure. The effect of para-tert-octyl phenolic (PTOP) resin on the tack of isoprene/N660 compound has been studied before and after ageing with FTIR spectroscopy [170]. The blends of XNBR and chlorobutyl rubber (CIIR) and XNBR-epichlorohydrin rubber (ECO) act as adhesive for aluminium-aluminium bonding [171,172]. In order to study the adhesion between CIIR-XNBR blend on aluminium foil, Bhattacharya and co-workers [173] used FTIR-ATR spectra [173], and they found that there was no change of spectra of the blend on Al-foil, i.e., on the leached aluminium surface, as indicated by the intensity of the carboxyl peak at 1736 cm-1 (for free COOH group or ester group) and 1696 cm-1 (for hydrogen bonded carbonyl). Manoj and co-workers [174] found that in addition to molecular interaction and physical entanglements of the molecular chains across the interface in PVC-NBR rubber joints, at high temperatures and long contact times, interfacial chemical bonds may be formed which seem to couple the two adherends, thereby resulting in cohesive failure of the rubber matrix on peeling. IR spectroscopic [174] studies of the PVC/NBR blend reveal formation of chemical bonds at the contact temperatures studied. The changes in the absorbance in the regions 3500-2900 cm-1 and 1750-1500 cm-1 correspond to the chemical interaction between the functional groups in the system are shown in Figure 3.12. The broadening of the peak at 3500-2900 cm-1 is assigned to N-H and O-H bonds formed by the hydrolysis of acrylonitrile (C≡N) group (2222 cm-1) in the presence of hydrogen chloride liberated during thermal degradation of PVC [131].

111


Figure 3.12 Infrared spectra of 50/50 (w/w) PVC-NBR blends moulded at 150 °C for 2 and 60 minutes Reprinted with permission N.R. Manoj and P.P. De, Journal of Adhesion, 1993, 43, 199. Copyright 1993, Gordon and Breach Science Publishers, [174]

112

Infrared Spectroscopy of Rubbers The increase in absorbance at 1730 cm-1 may be attributed to the formation of amide, acid or ester groups in the system. The amide and ester crosslinks are formed via the reaction between allylic chlorines in PVC and amide, acid groups in NBR [131]. The reduction in the peak at 1532 cm-1 may be due to the ring opening of the triazine derivatives in NBR, formed by the cyclisation of nitrile groups. The different amide bands could not be distinguished from the C=C stretching vibrations. The existence of an interfacial reaction between polyamide-6 and acrylate rubber (ACM) was confirmed by Jha and co-workers [112] with the help of IR spectra, in which joints made of ACM rubber and polyamide-6 annealed at 220 °C for different times indicated a reaction between the amine (-NH2) and carboxyl (-COOH) end groups of Nylon 6 with the reactive epoxy groups of ACM chains.

3.5.7 Degradation On storage, the stability of polymers changes due to their ageing as a result of the reaction of double bonds and other reactive groups. The degradation of ENR was studied with IR spectroscopy [175], which showed that during thermal ageing of ENR, carbonyl, alcohols, THF and ether, crosslinks are formed most probably due to the presence of a very small amount of acid. The thermal degradation of NR [176] in air for 1, 2 or 3 hours at 145 °C show the absorbance intensity of sample (37% gel content) decreases at 835, 1375, 1450, 1660 cm-1 with a prolonged heating whereas the absorbance at 1715 cm-1 and 3500 cm-1 increases, indicating the formation of aldehyde, ketone and carboxylic acid. FTIR is also applied for the development of biodegradable polymers suitable for biomedical application [177]. Two types of PU were made containing PEO (polyethylene oxide) and polycaprolactone (PCL), an amino acid based diester chain extender was used to confer degradability [177]. FTIR showed a large carbonyl peak (~1735 to 1724 cm-1) resulting from the ester groups of PCL-PU whereas the PEO polyurethanes exhibit a distinct ether peak at ª1100 cm-1. For PCL-PU the carbonyl stretching (1650-1800 cm-1) is dominated by intense soft-segment ester band, located at 1724 to 1735 cm-1 depending on the presence or absence of soft segment crystallinity [177].

3.6 Reverse Engineering Rubber industries produce various types of complicated products like tyre, cable, belt, seal, bearings, engine mounts, etc. The products are composed of rubber, plastics, fibre, metal, fillers and many other additives like antioxidants, accelerators, etc. Reverse engineering is a technique by which a rubber technologist can reconstruct the composition of the products based on thermal analysis. Dormagen [178] and Baranwal [179] reconstructed the formulation of a tyre, based on analyses of FTIR, spectra, thermal analysis and high performance liquid chromatography.

113


3.7 Conclusion It is concluded that IR spectroscopy provides information on qualitative as well quantitative analyses of rubbery materials, apart from their microstructures (that is, whether cis or trans, syndiotactic, atactic or isotactic). Different types of rubber blends (compatibilised or self-crosslinked) can be identified by the infrared spectroscopy. Synthesis, and degradation of polymers can also be followed by IR spectra. Mechanism of interaction between rubbers and fillers, can also be studied by IR-spectra. Different types of chemical reactions like the milling behaviour of rubbers, mechanism of adhesion and degradation can also be studied with the help of IR spectroscopy. The technique plays a great role in the product analysis under reverse engineering.

Acknowledgement Thanks are due to Mr. Rajeev R.S. and Mr. Shambhu Bhattacharyya for assistance in preparing the manuscript.

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Spectroscopy of Rubbers and Rubbery Materials 132. N.R. Manoj and P.P. De, Plastics and Rubber and Composites Processing and Applications, 1995, 23, 2, 103. 133. P. Ramesh and S.K. De, Polymer, 1993, 34, 23, 4893. 134. P. Ramesh and S.K. De, Polymer Networks and Blends, 1992, 2, 4, 209. 135. A. Mallick, D.K. Tripathy and S.K. De, Polymer Networks and Blends, 1993, 3, 1, 51. 136. T. Bhattacharya and S.K. De, European Polymer Journal, 1991, 27, 1065. 137. C.J. Paul, M.R.G. Nair, N.R. Neelakantan, P. Koshy, B.B. Idage and A.A. Bhelhekar, Polymer, 1998, 39, 26, 6861. 138. C.J. Paul, M.R.G. Nair, P. Koshy and B.B. Idage, Journal of Applied Polymer Science, 1999, 74, 706. 139. E. Yilgar, E. Burgaz, E. Yurtsever and I. Yilgor, Polymer, 2000, 41, 849. 140. L. Zha, M. Wu and J. Yang, Journal of Applied Polymer Science, 1999, 73, 2895. 141. C-T. Zhao and M.N. de Pinho, Polymer, 1999, 40, 6089. 142. Z. Wang, D. Gao, J. Yang and Y. Chen, Journal of Applied Polymer Science, 1999, 73, 2869. 143. R.N. Kumar, C.K. Woo and A. Abusamah, Journal of Applied Polymer Science, 1999, 73, 1569. 144. J.B. Donnet and A. Vidal, Advances in Polymer Science, 1986, 76, 103. 145. E.M. Dannenberg, Rubber Chemistry and Technology, 1986, 59, 512. 146. J.V. Hallum and H.V. Drushel, Journal of Physical Chemistry, 1958, 62, 110. 147. J.B. Donnet, Carbon, 1982, 20, 266. 148. E. Papirer, E. Guyon and N. Perol, Carbon, 1978, 16, 133. 149. C.H. Chen, J.L. Koenig, J.R. Shelton and E.A. Collins, Rubber Chemistry and Technology, 1982, 55, 103.

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Infrared Spectroscopy of Rubbers 150. A. Roychoudhury, Chemical Interaction of Chlorosulfonated Polyethylene with Functionalised Polymers and Surface Modified Fillers, IIT Kharagpur, India, 1994. [Ph.D. Thesis] 151. A. Roychoudhury and P.P. De, Journal of Applied Polymer Science, 1995, 55, 9. 152. A. Roychoudhury and P.P. De, Journal of Applied Polymer Science, 1993, 50, 181. 153. M.J. Wang, S. Wolf and J-B. Donnet, Rubber Chemistry and Technology, 1991, 64, 714. 154. A. Roychoudhury, S.K. De, P.P. De, J.A. Ayala and G.A. Joyce, Rubber Chemistry and Technology, 1994, 67, 662. 155. S. Bandyopadhyay, P.P. De, D.K. Tripathy and S.K. De, Journal of Applied of Polymer Science, 1995, 8, 719. 156. A.K. Manna, P.P. De, D.K. Tripathy and S.K. De, Rubber Chemistry and Technology, 1997, 70, 624. 157. S. Bandyopadhyay, P.P. De, D.K. Tripathy and S.K. De, Journal of Applied Polymer Science, 1997, 63, 1833. 158. Ajoy K. Manna, P.P. De, D.K. Tripathy, S.K. De and M.K. Chatterjee, Rubber Chemistry and Technology, 1999, 72, 2, 398. 159. M.J. Wang, S. Wolf and J-B. Donnet, Rubber Chemistry and Technology, 1991, 64, 559. 160. S. Bandyopadhyay, P.P. De, D.K. Tripathy and S.K. De, Plastics and Rubber and Composites Processing and Applications, 1996, 25, 7, 327. 161. A. Roychoudhury, P.P. De, N. Roychoudhury and A. Vidal, Rubber Chemistry and Technology, 1995, 68, 815. 162. M.L. Kralevich and J.L. Koenig, Rubber Chemistry and Technology, 1998, 71, 2, 300. 163. A.K. Manna, P.P. De, D.K. Tripathy, S.K. De and D.G. Peiffer, Journal of Applied Polymer Science, 1999, 74, 389.

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Application of Infrared Spectroscopy to Characterise Chemically Modified Rubbers …

4

Application of Infrared Spectroscopy to Characterise Chemically Modified Rubbers and Rubbery Materials Nikhil K. Singha and Prajna P. De

4.1 Introduction Chemical modification of polymers continues to be an active field of research [1-5]. It is a common means of changing and optimising the physical, mechanical and technological properties of polymers [5-7]. It is also a unique route to produce polymers with unusual chemical structure and composition that are otherwise inaccessible or very difficult to prepare by conventional polymerisation methods. For example, hydrogenated nitrile rubber (HNBR) which has a structure which resembles that of the copolymer ethylene and acrylonitrile, is very difficult to prepare by conventional copolymerisation of the monomers. Polyvinyl alcohol can only be prepared by hydrolysis of polyvinyl acetate. Most of the rubbers or rubbery materials have unsaturation in their main chain and/or in their pendent groups. So these materials are very susceptible towards chemical reactions compared to their saturated counterparts. Infrared spectroscopy is one of the most important tools used to characterise the chemical structure, composition and microstructure of different polymers [8-10]. In earlier chapters, the principles and applications of infrared (IR) spectroscopy in the characterisation of rubbers have been discussed. This chapter describes how IR spectroscopy can be used to characterise different types of chemically modified elastomers.

4.2 The Infrared Spectra of Commonly Used Diene Rubbers All diene rubbers may have few isomers depending on the addition of diene monomers. Different modes of addition are shown in Scheme 4.1. ‘R’ may be different groups, e.g., for chloroprene rubber (CR), R = Cl; for polyisoprene or natural rubber (NR) R = CH3).

125


Scheme 4.1 Microstructures in the diene elastomers

Butadiene rubber (BR) is the simplest case with R = H. It shows strong bands at 1625, 1643 and 1666 cm-1 for C=C stretching (n), for pendent -CH=CH2 (vinyl), cis 1, 4 BR and trans 1,4 BR, respectively. It also shows characteristic peaks at 967, 910 and 740732 cm-1 associated with wagging of the vinylic C-H bonds for trans 1,4, pendent and cis 1,4 double bonds, respectively [11]. However, these values may change for other diene rubbers which have analogous structure of BR, e.g., polyisoprene [12, 13] and CR [14] depending on the ‘R’ group in Scheme 4.1. According to Urey-Bradley field theory it is assumed that the important forces in molecular vibration act along chemical bonds and also includes interaction terms between adjacent nonbonded atoms. Such a force field can thus take into account attraction and repulsion between adjacent nonbonded atoms, which is known as field effects [15]. Slight deviation in different frequencies in some elastomers from those of BR may be ascribed to the field effect. The characteristic frequencies of different bands of various diene-based rubbers are shown in Table 4.1. It is interesting to note how the position and intensity of different absorption bands change during chemical modifications.

4.3 Hydrogenation Hydrogenation is an important method of chemical modification of elastomers. Because of the absence of carbon-carbon unsaturation, hydrogenated elastomers have good resistance to oxidative and thermal degradation, improved weatherability and good resistance towards chemicals and fluids [5-7]. Nitrile rubber (NBR) is a specialty rubber, and because of its oil resistance properties, it has been used in oil-wells and the automotive industry. Hydrogenation of NBR has been studied extensively because of its technological importance [16-19]. It was observed that on hydrogenation there was no change in the position or intensity of the peak at 2236 cm-1 (ν of -CN in acrylonitrile content) (Figure 4.1).

126


Table 4.1 Assignments of different characteristic bands in infrared spectroscopy of elastomers Assignments

Ref. [11]

967 910

C=C stretching of cis —CH=CH— C=C stretching of cis CH2=CH— CH2 deformation CH in-plane deformation of CH2=CHCH in-plane deformation of CH=CH— CH out-of-plane deformation (wagging) of CH2=CH— CH wagging of trans –CH=CHCH2 wagging of CH2=CH- unit

IR

1665 1645 1152 1130 890 845 840

C=C stretching of -CH=C(CH3)C=C stretching of CH2=C(CH3)C-CH3 of trans -CH=C(CH3)C-CH3 of cis -CH=C(CH3)CH wagging of CH2=C(CH3)CH wagging of trans CH=C(CH3)CH wagging of cis -CH=C(CH3)-

[12, 13]

CR

1660 1653 847 822 720

C=C stretching in trans CR C=C stretching in cis CR CH wagging in cis CR CH (of =CH-) wagging in trans CR —C-Cl stretching

[14]

NBR

2236 970 920 750

-CN stretching of acrylonitrile (ACN) CH wagging of trans -CH=CH- unit CH2 wagging of CH2=CH- unit CH wagging of cis -CH=CH- unit

[16-22]

SBR

1002 980 920 750 700

trans CH wagging of vinyl group CH wagging of trans -CH=CH- unit CH2 wagging of CH2=CH- unit CH wagging of cis -CH=CH- unit Characteristic for polystyrene

[27-32]

Elastomers BR

Characteristic Peaks (cm-1) 1660 1640 1470 1418 1408 995

127


Figure 4.1 Hydrogenation of NBR (34% acrylonitrile (ACN)) and hydrogenated NBR with 22.7, 11.0 and 0.6% mol unhydrogenated butadiene units Reprinted from [21] with permission from Huthig Publications, Heidelburg, Germany, Copyright 1989

The peaks at 970, 920 and 750 cm-1 which are due to ω(CH2) of 1,4 trans, ω(CH2) of pendent and ω(CH) of 1,4 cis content slowly disappear as the degree of hydrogenation increases. A new peak at 723 cm-1 appears and its intensity increases with increasing degree of hydrogenation. This peak is due to rocking of (CH2)n>4 in the hydrogenated NBR. Marshall and co-workers [20] and Bruck [21,22] have studied the infrared spectrum of HNBR (Figure 4.1). The degree of hydrogenation in HNBR can be calculated using the following equation: Degree of hydrogenation (mol%) = 100 - C(BR) x 100 C(BR) + C(HBR) where C(BR) is the relative amount of carbon-carbon unsaturation remaining in HNBR and C(HBR) is the relative amount of methylene group (CH 2) n formed during hydrogenation of olefin group in NBR. The molar concentrations of different components C(I) can be calculated by these equations,

128


=

C(ACN)

C(BR)

C(HBR)

=

1 ΣAi

A(970) 1 ⋅ K(970) ΣAi =

A(723) 1 ⋅ K(723) ΣAi

Where A is the normalised value of absorbances calculated using A(2236) as standard. K is the calibration factors for the absorption at different absorbance maxima. Bruck [21,22] calculated different ‘K’ values (K970, K723, K920, etc.), by using IR spectra as well as nuclear magnetic resonance (NMR) of different NBR and HNBR with different ACN content. Different spectral parameters of NBR and HNBR are shown in Table 4.2. Bhattacharjee and co-workers [23] hydrogenated liquid carboxylated NBR (XNBR) using homogeneous palladium and rhodium complexes. They observed that the absorption band at 1730 cm-1 (>C=O of free -COOH) and 1705 cm-1 (H bonded -COOH) remained unchanged when palladium catalysts were used. But the intensity of these peaks decreased on hydrogenation using a rhodium catalyst, which implied that there were substantial amounts of decarboxylation. The simplest diene rubber for hydrogenation is BR. The structure of the hydrogenated BR (HBR) depends on the different microstructures present in BR. Hydrogenation of BR with high 1,4 structure converts this elastomeric polymer into a tough semicrystalline

Table 4.2 Assignments of different peaks in IR spectra of NBR and HNBR Peak cm-1

Assignments

2236

-CN stretching of ACN

1.0

2280-2190

970

CH wagging of 1,4 trans (-CH=CH-)

2.30

1030-930

920

CH2 wagging of 1,2 unit (-CH=CH2-)

2.24

750

CH wagging of 1,4 cis (CH=CH)

0.36

860-660

723

(CH2)n rocking of HNBR (n >4)

0.25

840-680

Absorption Baseline drawn Reference factor(k) between point [21-23]

129

Spectroscopy of Rubbers and Rubbery Materials polymer which resembles the properties of low density polyethylene (LDPE) [24]. Hydrogenation of BR having predominantly 1,2 content leads to a polymer with thermoplastic elastomeric properties. The IR spectrum of cis-1,4 BR (98% cis C=C) and HBR is shown in Figure 4.2. The characteristic bands for the C=C unsaturation at 1640 cm-1 (C=C stretching for cisalkene), 1005 cm-1 (a small peak for trans CH wagging), 920 cm-1 (a weak peak for CH2 wagging for a vinyl group) and 750 cm-1 (a strong peak for CH wagging in cis-C=C) disappear confirming quantitative hydrogenation. Mohammadi and Rempel [25] hydrogenated BR containing 90% 1,2 units, which shows a strong peak at 910 cm-1 because of CH2 wagging of the 1,2 units. The disappearance of this peak and other characteristic bands for unsaturation (1640, 992 and 970 cm-1) signify quantitative hydrogenation in HBR. The study of hydrogenation using RhCl(PPh3)3 showed that 1,2 addition units were selectively hydrogenated over 1,4 addition units present in BR [24,25]. Hydrogenation of telechelic BR [26] was done either by diimide, a chemical generated

Figure 4.2 IR spectra of cis-1,4 BR and hydrogenated BR (HBR) Reprinted from X. Guo and G.L. Rempel, Journal of Molecular Catalysis, 1990, 63, 279, Figure 3, Copyright 1990, with permission from Elsevier Science

130

Application of Infrared Spectroscopy to Characterise Chemically Modified Rubbers … by a thermal decomposition of p-toluene sulfonyl hydrazide or in two steps via a partially saturated intermediate prepared by prehydrogenation up to 95% using gaseous hydrogen and a Ziegler-Natta catalytic system, followed by diimide hydrogenation up to 100%. It can be seen that after 100% hydrogenation, all bonds pertaining to C=C bonds (910, 968, 995 and 1640 cm-1) disappeared. Hydrogenation of the random and block copolymers of styrene and butadiene have been reported in the literature [27-32]. De Sarkar and co-workers [30-32] characterised hydrogenated styrene butadiene elastomer (HSBR) with IR spectroscopy. SBR shows bands at 699 cm-1 due to an aromatic unit, at 767 cm-1 due to CH2 wagging in 1,4 cis, unit at 909 cm-1 due to CH2 wagging in the 1,2 unit, and 967 cm-1 due to trans CH wagging in the butadiene content (Figure 4.3). As the degree of hydrogenation increases, the intensity of these peaks decreases and new peaks at 1450 cm-1 (-CH2 deformation in HSBR) and 723 cm-1 (-CH2 rocking vibration (CH2)n>4 in HSBR) appear in the spectra of HSBR (Figure 4.3a). The peaks due to aromatic styrene units are not changed, since the styrene units are not hydrogenated. So the peak at 699 cm-1 can be used for normalisation of other peaks in IR spectra of HSBR or partially hydrogenated HSBR. De Sarkar and co-workers [32] hydrogenated carboxylated SBR (XSBR containing 1-4% methacrylic acid) with diimide reduction technique in the latex stage using hydrazine hydrate, H2O2 and Cu2+ as catalyst. They observed selective hydrogenation of XSBR without affecting COOH groups. The infrared spectra of XSBR show a peak at 1694 cm-1 due to >C=O stretching vibration in COOH (Figure 4.3b). The peak shape and the frequency of >C=O absorption suggested that most of the COOH groups in XSBR were hydrogen bonded. The broad carbonyl peak for HXSBR can be deconvoluted into three peaks, one at original 1694 cm-1 and the other two approximately at 1720 and 1680 cm-1 (Figure 4.3b). The shift towards higher frequency was due to the reduction of strength of hydrogen bond in HXSBR whereas the peak shift towards lower frequency to 1680 cm-1 was ascribed to the ionic aggregates or clusters involving COOH groups and Cu2+ ions [32]. Polyisoprene can have few isomers according to different mode of addition as shown in Scheme 4.1. Cis 1,4 polyisoperene [12,13] shows a peak at 840 cm-1 whereas the trans isomer shows a peak at 845 cm-1 and 3,4 polyisoprene shows a characteristic peak at 890 cm-1. NR is cis 1,4 polyisoprene. Singha and co-workers [33] hydrogenated NR and observed that quantitative hydrogenation of NR led to an alternate copolymer of ethylene and propylene. The peaks at 840 and 1663 cm-1 due to the double bond disappeared [33] and a new peak at 735 cm-1 emerged due to (CH2)3 groups in hydrogenated NR [34] (HNR). Burfield [35] carried out the solid state hydrogenation of natural rubber in order

131


Figure 4.3a IR spectra of SBR and HSBR (70% and 94%) hydrogenation Reprinted from M. De Sarkar, P.P. De and A.K. Bhowmick, Journal of Applied Polymer Science, 1997, 66, 1151, Figure 5, with permission of John Wiley and Sons, Copyright 1997

to increase the ageing properties. Like NR, epoxidised natural rubber (ENR) was hydrogenated [36] in the presence of palladium acetate catalyst in a high pressure autoclave at a temperature of 323 K with hydrogen gas under a pressure of 2.7 MPa according to Scheme 4.2 [36]:

132


Figure 4.3b IR spectra of (a) XSBR and HXSBR with (b) 60% saturation (c) 80% saturation Reprinted from M. De Sarkar, P.P. De and A.K. Bhowmick, Polymer, 2000, 41, 907, Figure 3, Copyright 2000, with permission from Elsevier Science

133


Scheme 4.2 Hydrogenation of ENR Reprinted with permission from S. Roy, S. Bhattacharjee and B.R. Gupta, Journal of Applied Polymer Science, 1993, 49, 375. Copyright 1993, John Wiley and Sons, Inc.

The peak due to the 1,4 >C = C moiety, i.e., 1665 cm-1, is for reduced hydrogenated ENR (HENR). The peaks for 3,4 and 1,2 structures, at 1645 cm-1 and 1638 cm-1 disappeared in the case of HENR, signifying that they are more prone towards hydrogenation. The peak height of 1450 cm-1 corresponding to the >CH2 group increased, while the peak height at 870 cm-1 (for epoxy group) remained unchanged. CR may have four structural isomers, the trans 2-chloro-2-butenylidene-1,4, cis 2-chloro2-butenylidene-1,4 arising from 1,4 polymerisation and two other isomers arising from 1,2 and 3,4 polymerisation [14]. Cis and trans isomers show peaks at 1653 and 1660 cm-1 for C=C stretching, whereas the peaks due to CH wagging appear at 822 and 847 cm-1 for cis and trans isomers, respectively. Weak bands at 925 and 833 cm-1 appear due to CH wagging from 1,2 and 3,4 addition units in polychloroprene [14]. Singha and coworkers [37] hydrogenated CR using RhCl(PPh 3) 3 as a catalyst. They observed dehydrochlorination and hydrogenation in hydrogenated CR (HCR). The characteristic bands [14] at 1653, 1660, 925, 847, 833, 822 cm-1 disappeared and a new peak at 725 cm-1 emerged due to CH2 rocking which was similar to polyethylene. Hydrogenation of CR using RhCl(PPh3)3 yielded a product with a predominantly linear polyethylene sequence with minor proportions of vinyl chloride sequence [37].

134


4.4 Halogenation Halogenation of elastomers profoundly alters its adhesion and cure characteristics. Chlorination is carried out with either gaseous chlorine or sulfuryl chloride. The polymers are dissolved in chloroform or dispersed in a swollen state in chloroform. Chlorine gas is bubbled through the polymer solution or dispersion. Cis and trans 1,4 BR have been chlorinated in chloroform using chlorine gas [38, 39]. Figure 4.4 shows the infrared spectra of chlorinated cis and trans BR.

Figure 4.4 IR spectra of chlorinated cis and trans 1,4 polybutadiene Reprinted from N. Murayama and Y. Amagi, Journal of Polymer Science, Part B, 1966, 4, 119, Figure 2, Copyright 1966. Reprinted by permission of John Wiley & Sons, Inc.

135

Spectroscopy of Rubbers and Rubbery Materials On chlorination the strong band at 967 cm-1 in original trans BR which is due to C-H out of plane vibration of –C=CH for 1,4 trans configuration, weakens and slowly disappears as the chlorination proceeds. The band at 1090 cm-1 also disappears on chlorination. As the chlorination proceeds, the intensity of the band at 1250 cm-1, which corresponds to the band of polyvinyl chloride (PVC) at 1254 cm-1 of the C-H bending, gradually increases (Figure 4.4b). On the other hand, the bands at 724 cm-1 and 1643 cm-1 in cis BR, which are due to C-H wagging and C=C stretching, either decrease in intensity or disappear. A weak band around 965 cm-1 appears again, which indicates the conversion of cis isomer to trans isomer during chlorination [38, 39]. Chlorinated trans BR shows a peak at 650 cm-1 and chlorinated cis BR shows peaks at 680 cm-1 and 590 cm-1 which are due to the C-Cl stretching mode. According to Murayama and Amagi [38], during chlorination of trans and cis BR, addition takes place first and then substitution occurs. Chlorination of 1,4 BR leads to elastomers which resemble head to head PVC [38]. Bromination of BR can be carried out in n-heptane or chloroform with a solution of bromine in an organic solvent. On bromination the intensity of characteristic olefinic bands decrease, the peaks at 1235 cm-1 (due to CH2 bending in brominated BR) and at 1225 cm-1 (due to CHBr) increase [40]. Brominated BR shows bands at 785 and 550 cm-1 due to CBr stretching mode. Assignments of different peaks in chlorinated and brominated BR have been described in Table 4.3. Halogen can be introduced into elastomers by reactions with chlorocarbene. According to Konietzny and Biethan [41], dichlorocarbene can be added to BR to form elastomers with dicyclopropane rings. Dichlorocarbene was prepared in situ from chloroform and aqueous caustic soda using phase transfer catalysis. The modified elastomer showed a strong band at 810 cm-1 characteristic for CCl2 wagging vibration. Ramesan and Alex [42] modified SBR with dichlorocarbene prepared by alkaline hydrolysis of chloroform using cetyltrimethyl ammonium bromide as a phase transfer agent. Dicholoro carbene modified SBR showed improved resistance towards heat, flame and solvents. The modified SBR showed characteristic C-Cl absorption peaks at 806 and 1059 cm-1 which are due to cyclopropane ring [43]. It was observed that CCl2 initially reacted with the cis double bond. With the increase of reaction time, the intensity of absorption at 968 cm-1 (due to the trans isomer) and at 1653 cm-1 (due to the cis isomer) slowly decreased. The peak at 698 cm-1 remained unaffected implying that aromatic ring did not take part in the reaction. The absorption at 910 cm-1 (due to vinyl 1,2 content) remained almost unaffected during modification. Pinazzi and Levesque [44] reported modification of 1,4 polyisoprene with dichlorocarbene. As the reaction time increased, the intensity of the absorption band at 835 cm-1 decreased and a new peak emerged at 820 cm-1 due to C-Cl absorption and at 1025 cm-1 due to cyclopropane ring [43]. Okamoto and co-workers [45] reported γ-ray induced addition reaction of carbon tetrachloride (CCl4) to the vinyl group of syndiotactic 1,2 BR. They observed that addition

136


Table 4.3 Different characteristic bands in various chemical modification of elastomers Modification Hydrogenation

Characteristic Assignments Bands (cm-1)

-CH2- rocking of (-CH2-)n (n>4) -(CH2) in polyisoprene 3-

[16-37]

Chlorination

1250, 650 C-H bending in chlorinated BR, C-Cl 680 and 590 stretching in trans BR, C-Cl stretching in cis BR

[38, 39]

Bromination

1235 1225 C-H bending in brominated BR, C-H 785 and 550 bending in CHBr, C-Br stretching

[40]

Cyclisation

3040, 1175, 820

Hydrosilylation Hydroboration and hydroxylation Hydroformylation

Oxidation

Phosphonylation

725, 735

Ref.

CH stretching in trisubstituted double bond due to cyclisation in BR, due to terminal double bond in cyclised polyisoprene

[51, 52] [57, 58]

Due to –SiCH2-

[61]

3340

O-H stretching in hydroxylated BR

[64 ]

1050

C-O stretching in hydroxylated BR

2700

C-H stretching in –CHO

1725

>C=O stretching in -CHO

1410 and 1235

[65-67]

1260 and 1380

C-O-C ring vibration in epoxidised BR

[68-70]

1240

C-O-C ring vibration in epoxidised IR

[71 ]

1720

>C=O in polyketones

1710

>C=O in –COOH

3500

O-H stretching

2525-2725

-OH(P=O)OH

1150-1362

Free P=O

1087-1261

Hydrogen-bonded P=O

~ 800

[72, 73]

[74 ]

(P-O-)-C

137


Table 4.3 Continued Modification Sulfonation

Ionomer formation

Weathering and degradation

Characteristic Assignments Bands (cm-1)

1176

O=S=O stretching in –SO3H

881

S-OH stretching

150 to 400

Depending on cations in sulfonated ionomers

~1200

Asymmetric stretching of –SO3

1020

Symmetric stretching of –SO3

1550

>C=O stretching in carboxylate ionomers (depends on nature cation and ionic structure)

~3400

O-H stretching

1700-1720 1175 1642, 1533

>C=O stretching

Ref. [75-82 ]

[83-85] [87-90]

[109 115] [115]

Due to terminal >C= stretching in TPEE Due to –CONH2– of polyamide

of CCl4 to 1,2 BR is anti-Markovnikov and it leads to increase in molecular weight (Mn) due to crosslinking. On addition of CCl4, vinyl bands at 3080, 1840, 1640, 1415, 995 and 910 cm-1 decrease, CH2 bands at 2990-2860 and 1460 cm-1 increase with increase in molecular weight. Furthermore, intensity of the band at 680 cm-1 due to C-Cl stretching increases as the addition of CCl4 proceeds. Braun and co-workers [46] reported that reductive dechlorination of PVC using tri-n butyl tinhydride (n-Bu3SnH) leads to vinylchloride-ethylene copolymers. Copolymers were characterised by casting film from tetrahydrofuran (THF) solution in a potassium bromide disk. It is thought that the IR absorption peak at 750 cm-1 is ascribed to the (CH2)3 sequences and the peak at 720 cm-1 is due to the (CH2)n>5 sequences. As the dechlorination starts, the peak at 750 cm-1 due to CH2 sequences appears and intensifies with reaction time. The intensity of the peaks at 690 and 615 cm-1 due to the C-Cl stretching vibration slowly decreases. If the reduced PVC contains more than 46 wt% chlorine, only the absorption peak at 750 cm-1 appears in the IR spectra. If the chlorine content is less than 46 wt%, the peak at 720 cm-1 weakens. In this case the (CH2)n>5 sequences become more prominent than the -(CH2)3 –sequences [46].

138


4.5 Isomerisation In 1957, Golub [47] first reported the cis-trans isomerisation of an unsaturated macromolecule. He showed the successful transformation of the cis 1,4 BR into corresponding trans units by means of ultraviolet irradiation in the presence of a suitable sensitiser, which may be any one of a wide variety of organic bromine or sulfur compounds. In the photoisomerisation of BR or other diene elastomers, there is direct excitation of πelectrons of the double bonds to an antibonding state in which free rotation and hence geometrical interconversion can occur. When BR is irradiated in the presence of some sensitisers, such as allyl bromide or phenyl sulfide, the band at 710 cm-1 due to the cis isomer slowly weakens as the irradiation time increases. On the other hand, the intensity of the peak at 970 cm-1 which is due to the trans isomer slowly increases. From an initial cis/trans ratio of 57/43, BR was isomerised with allyl bromide in nitrogen to a structure with a cis/trans ratio of 10/90. In this case the band at 1450 cm-1 which is due to CH2 bending was taken as internal reference. No side reactions like cyclisation or degradation were observed in this isomerisation process. Under the similar reaction condition this isomerisation could not be induced in NR [47]. The methyl group attached to the double bond appears to have a stabilising influence on the given configuration. There are a variety of catalytic methods for cistrans isomerisation of 1,4 BR. They include heating the polymer with nitrogen dioxide or thiol acid or sulfur dioxide or some metal catalyst based on rhodium and iron. In this case the cis isomer forms a π-complex with the catalyst and subsequent release of the catalyst regenerates the more stable trans form. Golub and Rosenburg [48] reported photoinduced cis-trans isomerisation of the pendent 1,2 bonds in 1,2 – poly(cis-1,4 hexadiene) (CHD). On UV irradiation the bands at 700 cm-1 and 1666 cm-1 (due to -CH=CH-) in CHD weaken and the intensity of the peak at 961 cm-1 due to the trans isomer increases.

4.6 Cyclisation Cyclisation of unsaturated elastomers has been the subject of interesting research. Cyclisation yields hard resinous products which have commercial importance and are designated as cyclised rubber [49, 50]. Cyclisation can be carried out with cationic, radiation induced, photo-induced or by thermal methods. Among all these methods cationic cyclisation has been extensively reported probably because there are less side reactions. The generally accepted mechanism of cyclisation of 1,3 diene elastomers is shown in Scheme 4.3a.

139


Scheme 4.3a Cyclisation of diene elastomers

Scheme 4.3b Monocyclic structure of 1,4 polyisoprene Reprinted with permission from M. Stolka, J. Vodehnal and I. Kossler, Journal of Polymer Science, Part A, 1964, 2, 3987. Copyright 1964, John Wiley & Sons, Inc.

The number of rings in the cyclised polymers depends on conditions of the cyclisation and on the nature of the polymers. Cis 1,4 BR has been cyclised with H2SO4 [51], TiCl4 [51] and alkyl aluminium halide in combination with an organic halide [52]. Cationic cyclisation of BR is more difficult with respect to polyisoprene rubber (IR) because of the lower stability of the secondary carbonium ion in BR compared to the highly stable tertiary carbonium ion in IR. Presence of the less stable carbonium ion in BR makes it difficult to control the side reactions and hence either chain scission or crosslinking occurs. Analysis of the infrared spectrum shows that the bands characteristic of cis 1,4 polybutadiene at 740, 1315 and 3008 cm-1 (cis 1,4 units), 910 and 990 cm-1 (1,2 units) and 967 cm-1 (trans 1,4 units) disappear and two new peaks emerge at 820 and 3040 cm-1. The peak at 820 cm-1 can be ascribed to out-of-plane deformation vibration and the peak at 3040 cm-1 is due to valency vibrations of the tri-substituted double bond, C=CH-. These bonds are created in the ring when BR undergoes intramolecular cyclisation. Figure 4.5 shows IR spectra in the 700-1100 cm-1 region of BR with different degrees of unsaturation. It shows that the intensity of the band at 740 cm-1 (cis 1,4 BR) falls off sharply and a new peak appears at 820 cm-1 whose intensity increases, as the degree of cyclisation increases.

140


Figure 4.5 IR spectra of cis 1,4-polybutadiene with (1) 97% unsaturation; (2) cyclised polybutadiene with 58%; (3) 36% and (4) 30% unsaturation Reprinted with permission from V.S. Shagov, A.I. Yakubchik and V.N. Podosok, Orskaya, Polymer Science USSR, 1969, 10, 1092. Copyright Pergamon Press, UK, 1969

The intensity of the band at 967 cm-1 (due to the trans 1,4 unit) increases a little in the initial stage and then it decreases. It indicates that in addition to cyclisation, cis-trans isomerisation occurs at the initial reaction stage. Figure 4.5 shows a fully cyclised BR (unsaturation 30%) though there is still some absorption at 740 cm-1 and 967 cm-1. Although it is in fully cyclised BR, there is still a small proportion of segments with double bonds, which cannot take part in further cyclisation because the neighbouring double bonds have already reacted. Information about the amount and the structure of the cyclic units present in cyclised BR (CBR) can be obtained by different spectroscopic methods. C13 NMR yields the quantitative information about unsaturation in CBR, i.e., linear and cyclic double bonds. IR spectra provide information on the residual linear unsaturation. According to the cyclisation mechanism it is assumed that one polycyclic sequence gives rise to one unsaturation. The fraction of polycyclic units and the average number of cyclohexane

141

Spectroscopy of Rubbers and Rubbery Materials rings present in each polycyclic sequence can be calculated using the equation explained in [53-56]. Priola and co-workers [52, 53] reported the average length of polycyclic sequences to be between 2.5 and 4 depending on the extent of cyclisation in BR. Quantitative IR and NMR spectra confirm the presence of the cyclic structures in some polymers [54]. Cyclisation of cis-1,4 (Hevea), trans-1,4 (balata) and 3,4 polyisoprene has also been reported using H2SO4, TiCl4 and p-toluenesulfonic acid as catalysts [57, 58]. Stolka and co-workers [57] studied cyclisation of different isomers of polyisoprene using H2SO4 and characterised them with IR spectroscopy. The intensity of the bands at 1315 cm-1 due to C=C-H vibration in –C(CH3)=CH- groups in cis form and at 1130 cm-1 assigned to the same group decreased in 1% H2SO4 and disappeared in concentrated H2SO4. The absorption of CH2 groups near 1455 cm-1 shifted to 1465 cm-1. When the cyclisation was complete, a very broad and strong band appeared at 1175 cm-1. Cyclisation leads to monocyclic, bicyclic or polycyclic structures. The monocyclic structure is shown in Scheme 4.3b. The peak at 1175 cm-1 is due to the terminal double bonds. It has been confirmed by studying the IR spectra of model cyclopolymers [58, 59]. In the completely cyclised state, new peaks appeared at 810, 885, 1040, 1175 and 1265 cm-1 together with a very weak band at 1200 cm-1. In the range of C-H stretching vibrations around 3000 cm-1 the absorption intensity at 2750 and 3040 cm-1 decreased and new bands appeared at 2685 and 3060 cm-1. During cyclisation the characteristic peaks of trans 1,4 polyisoperene at 1330 and 1150 cm-1 slowly decreased and then disappeared in the fully cyclised polymer. As in the case of cyclised cis isomer, a new peak emerged at 1175 cm-1. It was further observed that cyclised 1,4 cis and 1,4 trans isomers had identical IR spectra in the whole range 400 to 4000 cm-1.

4.7 Hydrosilylation Hydrosilylation is also a very useful chemical modification which leads to silane modified polymers with special properties [60-62]. Silane modified polymers have improved adhesion to fillers and better heat resistance. It also acts as a reactive substrate for grafting or moisture catalysed room temperature vulcanisation. Guo and co-workers [61] carried out catalytic hydrosilylation of BR using RhCl(PPh3)3 as the catalyst. Hydrosilylation reactions followed anti-Markovnikov rule as shown in the Scheme 4.4. The infrared spectra of BR and hydrosilylated BR are shown in Figure 4.6. As the degree of hydrosilylation increases, the absorbance at 3100 and 910 cm-1 due to the terminal C=C bonds slowly disappear. There are new absorbances at 1235 cm-1 due to CH2 scissoring for the SiCH2 structure in the hydrosilylated polymer. Hydrosilylation 142


Scheme 4.4 Hydrosilylation in diene elastomers Reprinted with permission from X. Guo and G.L. Rempel, Macromolecules, 1990, 23, 5047. Copyright 1990, American Chemical Society

Figure 4.6 (a) IR spectrum for PBD (A); (b) IR spectrum for PBD(A) + His (C2H5)3 (1) product Reprinted with permission from X. Guo, R. Farwaha and G.L. Rempel, Macromolecules, 1990, 23, 5047, Figure 2. Copyright 1990, American Chemical Society

143

Spectroscopy of Rubbers and Rubbery Materials of SBR proceeds with greater difficulty than that of BR. Presence of the aromatic ring in the copolymer chain creates a different electronic environment as well as the steric environment for the reaction centre in the polymer chain. With hydrosilylation the peaks at 915 cm-1 (=CH2 wagging) decreased and new peaks at 1220 cm-1 and 1005 cm-1 emerged due to newly introduced –SiCH2- structure. Hydrosilylation of NBR leads to Markovnikov addition products. Presence of -CN groups in NBR induces different electronic environment [62].

4.8 Hydroboration The hydroborated polymers are valuable intermediates that can be converted to a variety of functional polymers [63, 64]. Chung and co-workers [64] carried out hydroboration of 1,2 BR, 1,4 BR and 1,4 polyisoprene using 9-borabicyclo [3,3,1] nonane (9BBN) and then hydroxylated the polymers with NaOH/H2O2. The relative reactivity of the various polydiene is 1,2 BR>1,4 BR> 1,4 polyisoprene in the ratio 1000:167:125. As the degree of hydroboration and subsequent hydroxylation increase, all absorption peaks for vinyl groups at 3065, 1635, 990 and 905 cm-1 decrease. Two strong absorption bands emerge at 3340 cm-1 and 1050 cm-1 due to ν(OH) and ν(CO), respectively, in hydroxylated BR.

4.9 Hydroformylation A hydroformylation reaction in diene polymers introduces a formyl group which is an extremely reactive functional group. Sibtain and Rempel [65] carried out hydroformylation of SBR using HRh(CO)(PPh3)3 and reported anti-Markovnikov addition product. Hydroformylation takes place preferentially in the 1,2 unit. As the degree of hydroformylation increases new absorption bands appear at 1724 cm-1 due to ν(C=O) and at 2700 cm-1 due to ν(C-H) in CHO. Bhattacharjee and co-workers [66] carried out hydroformylation of NBR and observed new peaks at 1724 and 2700 cm-1 which are characteristics of CHO groups. Sanui and co-workers [67] carried out hydroformylation of polypentenamer (PPA). They converted hydroformyl group of the modified PPA to the aldoxime and subsequently to nitrile derivatives (Scheme 4.5). Then they carried out hydrogenation to convert the amorphous, unsaturated nitrile derivative to the crystalline saturated polymer. Figure 4.7 and Scheme 4.5 show the sequential chemical modification of PPA to nitrile containing hydrogenated PPA. Disappearance of peaks at 1724 cm-1 and 2700 cm-1 in PPA-CN indicates complete transformation of CHO, while emergence of the peak at 2240 cm-1 indicates presence of CN groups in PPA-CN.

144


Scheme 4.5 Chemical modification in polypentenamer (PPA) Reprinted with permission from K. Sanui, W.J. MacKnight and R.W. Lenz., Macromolecules, 1974, 7, 952. Copyright 1974, American Chemical Society

Figure 4.7 IR spectra of PPA-CHO, PPA-CN and HyPPA-CN Reprinted with permission from K. Sanui, W.J. MacKnight and R.W. Lenz, Macromolecules, 1974, 7, 952, Figure 1. Copyright 1990, American Chemical Society

145


4.10 Oxidation Oxidation induces polarity in the diene elastomers and hence increases resistance to hydrocarbon solvent. Oxidation of diene elastomers with peracids leads to epoxidation [68]. There are two types of epoxy rings depending on the chain microstructures. Epoxy groups may be in the main chain (internal groups) or in the pendent groups (external groups) [68]. Epoxidised cis-1,4 BR [69] shows characteristic peaks at 1260 and at 1380 cm-1 for ring stretching vibration (ν COC) and at 800 and 885 cm-1 for ring vibration. Epoxidised trans 1,4 BR shows characteristic peaks at 1060 and 1365 cm-1 for ring stretching vibration (ν COC) and at 710 and 880 cm-1 for ring vibration. Zuchowska [68] showed that during epoxidation some secondary reactions like acid catalysed ring opening take place. He observed a very intense peak at 1740 cm-1 due to hydroxyacetyl group in the spectra of 1,2 BR. Epoxidised cis 1,4 polyisoperene [69] showed peaks at 1240 cm-1 for νCOC and other characteristic peaks at 680 and 1070 cm-1. With epoxidation the intensity of the absorption bands due to carbon-carbon double bond decreases. Roy and co-workers [70] studied the epoxidation of NR with performic acid in the latex stage and their study reveals that the furanisation process (peak at 1068 cm-1) predominates only at epoxy contents above 65%. Diene elastomers can be oxidised to polyketones by a catalytic method [71]. Polyketones have potential uses, as they are good precursors for the synthesis of many other polymers. Polyketones are easily photodegradable. Iraqi and Cole-Hamilton [71] oxidised polybutadiene to polyketones with t-BuOOH in the presence of [Pt (diphoe) CF3(CH2Cl2)]+ (diphoe stands for Ph2PCH =CHPh2). IR spectra of all the products showed strong absorption at 1720 cm-1 indicating the induction of >C=O group in the polymer backbone. They achieved up to 57% conversion of C=C to >C=O. The reactivity of different microstructures follow the order [71], pendent BR > trans BR > cis BR. Cole-Hamilton and co-workers [72, 73] carried out catalytic hydrocarboxylation of BR using [PdCl2(PPh3)2]SnCl2 as the catalysts. Infrared spectra showed the disappearance or diminution of peaks at 732-740, 910, 967, 1650 cm-1 which are characteristic of carboncarbon unsaturation. A strong peak at 1710 cm-1 (ν C=O) and a very broad new peak (ν O-H) emerged due to the presence of the COOH group. Up to 93% of the unsaturation can be converted to COOH groups. The hydrocarboxylated BR was found to be useful in wood preservation as the polymer prevents the uptake of water [73].

4.11 Phosphonylation It is difficult to introduce phosphorus into a linear polymer by conventional polymerisation or co-polymerisation of phosphorus containing monomers. The polymers containing

146

Application of Infrared Spectroscopy to Characterise Chemically Modified Rubbers … phosphorus shows higher thermal stability and have better flame-retardant properties. A phosphate salt containing polymer shows microphase separation of the salt groups to form ionic domains. Azuma and Macknight [74] introduced dimethyl phosphonate ester groups into PP using benzoylperoxide (BPO) as free radical initiator. They achieved 10% phosphonylation of PP and hydrolysed using hydrochloric acid and then prepared its caesium salt using CsOH. Phosphonylated PP shows a peak at 820 cm-1 which is absent in its acid or salt form indicating it is due to P-O-C vibration. The different absorption bands in phosphonylated PP and its assignment have been shown in the Table 4.3.

4.12 Sulfonation Sulfonation is very useful chemical modification of polymer, as it induces high polarity in the polymer changing its chemical as well as physical properties. Sulfonated polymers are also important precursors for ionomer formation [75]. There are reports of sulfonation of ethylene-propylene diene terpolymer (EPDM) [76, 77], polyarylene-ether-sulfone [78], polyaromatic ether ketone [79], polyether ether ketone (PEEK) [80], styrene-ethylenebutylene-styrene block copolymer, (SEBS) [81]. Poly [bis(3-methyl phenoxy) phosphozene] [82]. Sulfonated polymers show a distinct peak at 1176 cm-1 due to stretching vibration of O=S=O in the -SO3H group. Another peak appears at 881 cm-1 due to stretching vibration of S-OH bond. However, the position of different vibrational bands due to sulfonation depends on the nature of the cations as well as types of solvents [75, 76].

4.13 Ionomer Formation Ionomers are polymers which contain up to about 10% mole percent of ionic group [8385]. Ionomers show interesting properties because of the presence of different interactions which include hydrogen bonding, formation of charge transfer complexes and ion-ion interactions. Polymers with carboxylic or sulfonic acid groups in its backbone on neutralisation with zinc, sodium or other metal salts form ionic aggregates. IR spectroscopy can be used to characterise the formation of ionomers by studying the environment of the anions [85, 86]. Risen and co-workers [87, 88] used far-IR spectra (150 to 400 cm-1) to demonstrate the sensitivity of low frequency vibrations to the anions and cations and the degree of cluster formation in ionomers. For example, styrene sulfonic acid ionomers with Na+ cation shows absorption bands at 220 cm-1, whereas the Cs+ cation shows bands at 100 cm-1. Carboxylate ionomers have been characterised with Fourier transform-infrared (FT-IR) in the region of antisymmetric stretching vibration of carboxylate anions. Figure 4.8 shows carboxylate ionomer [89] of ethylene methacrylic (4%) copolymer).

147


Figure 4.8 FT-IR spectra of carboxylate ionomers ionised with zinc salt in the range of 1900-1200 cm-1 Reprinted with permission from B.A. Brozoski, M.M. Coleman and P.C. Painter, Macromolecules 1984, 17, 230, Figure 4. Copyright 1990, American Chemical Society

The absence of peaks at 1750-1700 cm-1 indicates complete ionisation. The nature of band depends on the types of cations. The salts of potassium, caesium and zinc exhibit sharp single bands whereas sodium, calcium, strontium and barium salts show distinct doublet bands [90]. It is due to different ionic environment with various cations.

148

Application of Infrared Spectroscopy to Characterise Chemically Modified Rubbers … Sulfonated ionomers are also characterised by IR spectroscopy [75-77]. Agarwal and coworkers [76] analysed the Zn+2 salt of sulfonated EPDM. The peak at about 1200 cm-1 is due to the asymmetric stretching of the sulfonate group. The band at 1020 cm-1 is ascribed to the symmetric stretching of the -SO3- group. The position and nature of the absorption band depend on the nature of the cation [76]. The band at 610-615 cm-1 is due to C-S stretching of the polymer -SO3- band. Mixed anionic (sulfonated – carboxylated) ionomers [81] were prepared by sulfonation of maleated block-copoly (styrene/ethylene-butylene/styrene) (m-SEBS) by acetyl sulfate, followed by neutralisation of the sulfonated maleated product, leading to the formation of a new block copolymer ionomer based on both carboxylate and sulfonate anions according to Scheme 4.6. FT-IR spectra confirm the presence of both carboxylated and sulfonate ions (Figure 4.9).

Figure 4.9 FT-IR spectra of m-SEBS, Na-m-SEBS and Na-s-m-SEBS Reprinted from S.K. Ghosh, D. Khastgir, S.K. De, P.P. De, R.J. Albalak and R.E. Cohen, Plastics, Rubber and Composite Processing and Applications, 1998, 27, 310, Figure 1, with permission from the Institute of Materials, UK, Copyright 1998

149


Scheme 4.6 Sulfonation reaction of m-SEBS Reprinted with permission from S.K. Ghosh, D. Khastgir, S.K. De, P.P. De, R.J. Albalak and R.E. Cohen. Plastics, Rubber and Composite Processing and Applications, 1998, 27, 310. Copyright 1998,

150

Application of Infrared Spectroscopy to Characterise Chemically Modified Rubbers … The characteristic absorbance of m-SEBS occurring at 1716 cm-1 indicates the C=O stretching frequency of a five membered maleic anhydride group in the hydrogen bonded maleic acid form. The absorbance at 700 cm-1 is due to the C-H rocking vibration of the polystyrene ring and the 1602 cm-1 band is due to phenyl ring vibration of polystyrene. The carboxylated salt (Na-m-SEBS) shows the occurrence of two new bands at 1576 cm-1 and 1548 cm-1 as a doublet which is ascribed to the asymmetric carboxylate stretching vibration [90]. The Na-s-m-SEBS shows an absorbance at 1048 cm-1, which is due to the symmetric stretching of the SO3 - anion of the sodium sulfonate group [76]. Other characteristic bands at 1581 and 1552 cm-1 are due to a doublet resulting from the octahedral arrangement of the carboxylate anions. Ghosh and co-workers studied such mixed anionic ionomers with zinc [91], magnesium [92] and barium [92].

4.14 Ionomeric Blends The ionic aggregates present in an ionomer act as physical crosslinks and drastically change the polymer properties. The blending of two ionomers enhances the compatibility via ion-ion interaction. The compatibilisation of polymer blends by specific ion-dipole and ion-ion interactions has recently received wide attention [93-96]. FT-IR spectroscopy is a powerful technique for investigating such specific interactions [97-99] in an ionic blend made from the acid form of sulfonated polystyrene and poly[(ethyl acrylate – CO (4, vinyl pyridine)]. Datta and co-workers [98] characterised blends of zinc oxideneutralised maleated EPDM (m-EPDM) and zinc salt of an ethylene-methacrylic acid copolymer (Zn-EMA), wherein Zn-EMA content does not exceed 50% by weight. The blend behaves as an ionic thermoplastic elastomer (ITPE). Blends (Z0, Z5 and Z10) were prepared according to the following formulations [98]: Z0 m-EPDM 100, ZnO 10 Z5 Zn-EMA 50, m-EPDM 50, ZnO 10 Z10 Zn-EMA 100, ZnO 10 A typical infrared spectra of 50/50 blend of m-EPDM and Zn-EMA [98] is shown in Figure 4.10. The m-EPDM-ZnO system shows a broad diffused band at 1565 cm-1 which is believed to be due to asymmetric stretching of bridging type carboxylate groups. The Zn-EMA-ZnO system on the other hand, shows a high intensity band at 1586 cm-1 indicating a high extent of metal carboxylate salt formation. The asymmetric stretching peak in the case of mix Z10 occurs at higher frequency than that in mix Z0, due to the occurrence of stronger ionic aggregates or clusters in mix Z10. In the absence of any interaction, the spectrum of the blend is likely to be equivalent to the summation spectrum (Figure 4.10c) of the 151


Figure 4.10 IR spectra of (a) Z0 (m-EPDM); (b) Z10 (Zn-EMA); (c) Summation averaged spectrum; (d) Z5 (blend 50/50 m-EPDM and Zn-EMA) Reprinted with permission from S. Datta, P.P. De and S.K. De, Journal of Applied Polymer Science, 1996, 61, 1839. Copyright 1996, John Wiley and Sons, Inc.

152

Application of Infrared Spectroscopy to Characterise Chemically Modified Rubbers … constituent polymers [98], indicating the absence of any interaction, but the blend corresponding to mix Z5 (Figure 4.10d) shows only a sharp peak at 1586 cm-1. It is expected that if the two ionic peaks combine, the intermediate band position should be around 1574 cm-1, but the band has been shifted to higher frequency side by 12 cm-1, which may be due to intermolecular interaction. Recently De and his co-workers have characterised a series of such ionomeric polyblends with infrared spectroscopy, as given in Table 4.4.

4.15 Weathering and Degradation of Polymers It is possible to study the effect of weather on elastomers by IR spectroscopy. However, IR bands of polymers are inherently broad and weak and it is difficult to detect minor chemical changes on the polymer chain. But it is possible to recognise these differences by using absorption subtractions of the original polymer from the reacted polymer. Koenig [109] used the absorbance subtraction method to study the surface oxidation of BR at 30 °C. When BR was exposed to weather for 10 hours there was a change in the ratio of cis and trans unsaturation and occurrence of a peak at 1065 cm-1 (for C-O) suggested formation of an oxidised product. After longer exposure oxidation gave rise to absorptions for OH (3300 cm-1) and C=O (1700 cm-1, 1720cm-1). Irradiation damage of polyethylene was studied in the same way [110]. Chlorosulfonated polyethylene (CSM) is a specialty elastomer because of its interesting properties. The changes in CSM on γ-irradiation have already been studied [111]. The most important changes occur in the range of 3600-3200 cm-1 and 1800-1650 cm-1, where the accumulation of OOH/OH and C=O groups occur. On γ-irradiation new peaks appear at 3400 cm-1 and 1725 cm-1 which are due to OH stretching and >C=O stretching, respectively. The nature of chemical changes depends on the intensity of dose of γ-radiation. Even after exposure at relatively small dose of radiation there are modifications of the bands at 1370 and 1155 cm-1 due to SO2Cl and SO2. The quantitative change of SO2Cl can be detected by the absorbance ratio at the band 1370 and 1460 cm-1 (due to C-H bending in PE). On γ- irradiation there is decrease in the intensity of the peak at 1370 and 1155 cm-1 due to SO2 evolution. New peaks emerge at 1610 and 825 cm-1 due to C=C unsaturation generated during modification. Weathering of acrylonitrile-butadiene-styrene terpolymer (ABS) has been characterised by IR spectroscopy [112] and it was shown that the butadiene unit is responsible for UV degradation. Weatherability of thermoplastic polyester elastomers (TPEE) based on polybutylene terephthalate and polytetramethylene glycol was studied using outdoor exposure [113]. The IR spectra of TPEE exposed to air for 6 months for the original TPEE were compared. New peaks occurred in exposed TPEE at 1175 cm-1 (due to >C= stretching). The ratio of the peaks at 1175 cm-1 to the absorbance at 1500 cm-1 was due to in-plane skeletal vibration of the benzene ring which determines the extent of weathering in the sample.

153


Table 4.4 Characteristic bands in ionomeric polyblends Constituents of polyblend 1. Zn-s-EPDM and Zn EMA

2. Zn-m-EPDM and Zn-m-HDPE

Main peaks Peak assignment in cm-1 1694

>C=O stretching

1584

bridging metal carboxylate

1259

coupled vibration of C-O stretch and OH bending

1562

asym. metal carboxylate stretching

1462

-CH2- bending

1366

-CH2-wagging

Ref. [100 ]

[101 ]

1538, 1563 asym. carboxylate stretching pair, octahedral carboxylate stretching 1596 3. Zn-m-HDPE and Zn-XNBR carboxylate

1596,1552

a doublet in Zn-m-HDPE, asym. stretching of carboxylate

a doublet due to rocking vibration in Zn-mHDPE

1570, 1536 a doublet carboxylate asym. stretching in octahedral metal carboxylate in Al-m-EPDM 561

Al stearate in Al-CSM

1016

symm. stretching of sulfonate group

1156

asym. stretching of S=O

5. Zn-m-HDPE 1894, 1775 C=O stretching of five membered maleic and Zn-m- EPDM anhydride 1709

hydrogen bonded carboxylic acid pairs

1376

CH3 symm. Deformation

1562

assym. carboxylate stretching, of Ionomer

1588, asymm. carboxylate stretching 1542, 1532

154

[102]

1587, 1541 asym. stretching of carboxylate in Zn-XNBR 720, 731

4. Al-m-EPDM Al-CSM

tetrahedral carboxylate vibration

[103]

[104]


Table 4.4 Continued Constituents of polyblend 6. Zn-m-EPDM and Zn-PEA

7. Zn-PPA and Zn-XNBR

8. Zn-m-EPDM and Zn-PEA

Main peaks Peak assignment in cm-1 1709

Hydrogen bonded carboxy lic acid pairs

1464

—CH2 bending

1376

symm. Deformation

1560

asymm. carboxylate stretching in ionomer

1534

asymm. stretching of carboxylate ion

1547, 1588 asymm. stretching of metal carboxylate 1458

—CH2 bending

1375

CH3 asymm. Deformation

1730

free carboxylic acid

1697

hydrogen bonded carboxylic acid

1448

coupled vibration of—CH2 – bending and –CH3 asymm. deformation

1714

hydrogen bonded carboxylic acid

1464

CH2 bending

1363

CH2 wagging

1560

assym. stretching of carboxylate

1630

acid salt complex

1595

asymm. tetrahedral carboxylate stretching

Ref. [105]

[106]

[107]

1564, 1541 asymm. octahedral zinc-carboxylate stretching 9. Zn-XNBR and Zn-PEA

1665

C=C stretch

[108]

1591, 1538 asymm. stretching of zinc carboxylate ion 1464

—CH2- bending

1714

hydrogden bonded carboxylic acid

1590

tetrahedral structure of zinc carboxylate ion

1562/1546 1620

octahedral structure of zinc carboxylate acid salt complex

155

Spectroscopy of Rubbers and Rubbery Materials Bhattacharjee and co-workers [114] studied degradation of NBR and HNBR using IR spectroscopy and X-ray photoelectron spectroscopy (XPS). They observed >C=O and -COOR functionalities in HNBR and NBR after ageing. The extent of ageing was quantified from the ratio of absorbances at 1735 cm-1 and 1463 cm-1 (due to CH2 deformation in HNBR) and 1732 cm-1 and 1446 cm-1 (due to CH2 deformation in NBR). Ghosh and co-workers [115] studied the thermal degradation of segmented polyamides at different temperatures with the help of IR spectra. In segmented polyamide [115], the peak at 1727 cm-1 is indicative of the ester linkage connecting the soft polyether segment and the polyamide hard block. The peaks occurring at 3311, 1642, 1553 cm-1 are due to the presence of –CONH2- groups, whereas the peaks at 1036, 1116, 1185 and 1275 cm-1 denote the various types of deformations taking place in the –CH2-O-CH2- linkages of the soft polyether segments. IR spectra of the degraded sample in air at 367 °C reveal the presence of major peaks at 3311, 1642 and 1533 cm-1. The absence of peaks at 1727, 1275, 1185 and 1036 cm-1 suggests that the polyether soft blocks are adversely affected by oxidative degradation.

4.16 Modification by Radiation Modification of polymers by radiation is a potential method for the development of new polymers and composites. Radiation crosslinking thus endows the polymers with special properties. Datta and co-workers used IR spectroscopy to characterise modified ethylene vinyl acetate (EVA) exposed to electron beam in the presence of trimethylolpropane trimethacrylate (TMPTMA) [116] or triallyl cyanurate (TAC) [117]. The IR spectra of pure EVA shows peaks at 1740 cm-1 (>C=0 of ester group), 1480 cm-1 (-CH2 blending), 1370 cm-1 (-C-H bending of CH3), 1260 cm-1 (C-O stretching) and 1030 cm-1 (-C-O-C, ether linkage). The spectra of unirradiated blend of EVA and TMPTMA show a sharp peak at 1640 cm-1 mainly resulting from >C=C stretching of trans-vinylene, present in TMPTMA. When the blends are irradiated with 50 KGy doses of electron beam, the peak at 1640 cm-1 disappears, due to grafting or crosslinking of TMPTMA with EVA through unsaturation. Up to 150 KGy dose, various reactions like aerial oxidation, crosslinking of EVA with TMPTMA, chain scission, disproportionation and cyclisation reactions take place [116]. IR study showed that some residual unsaturations remained in irradiated pure TMPTMA, while in blends all unsaturations were used up at a very early stage of irradiation. Some ether linkages were formed during irradiation in pure EVA and the blends [118], although in pure EVA the concentration of ether linkages reached a maximum at the 2 Mrad dose and then decreased, while in blends it increased with an increase in radiation dose [118]. Electron beam initiated grafting of TMPTA onto EPDM has been carried out over a wide range of irradiation doses (0-200 kGy) using a fixed concentration (10%) of TMPTA [119, 120]. IR studies indicate increased peak absorbances at 1730, 1260, 1120 and 1019 cm-1 up to 50 kGy and hence increased

156

Application of Infrared Spectroscopy to Characterise Chemically Modified Rubbers … >C=O and C-O-C concentrations, which may cause crosslinking as well as chain scission of the EPDM rubber. On the other hand, free radicals developed on trimethylol propane triacrylate (TMPTA) participate in reactions like self-crosslinking, grafting, cyclisation and copolymerisation. Similarly, the influence of polyfunctional monomers such as tripropylene glycol diacrylate (TPGDA), TMPTA, TMPTMA, TAC on the structural changes of fluorocarbon rubber (Viton B) exposed to electron beam has been investigated with the help of attenuated total reflectance-IR spectroscopy [121]. The absorbance at 1397, 1021, 672, 504 cm-1 due to C-F group decrease on irradiation of the mixtures of fluoroelastomer and TMPTA, indicating dehydrofluorination and scission. The photopolymerisation of tetra chloroethyl acrylate, pentafluorophenyl acrylate, pentafluorophenyl-methacrylate was examined by quantitative FT-IR [122].

4.17 Conclusion IR spectroscopy can be used to characterise not only different rubbers, but also to understand the structural changes due to the chemical modification of the rubbers. The chemical methods normally used to modify rubbers include hydrogenation, halogenation, hydrosilylation, phosphonylation and sulfonation. The effects of oxidation, weathering and radiation on the polymer structure can be studied with the help of infrared spectroscopy. Formation of ionic polymers and ionomeric polyblends behaving as thermoplastic elastomers can be followed by this method. Infrared spectroscopy in conjunction with other techniques is an important tool to characterise polymeric materials.

Acknowledgement NKS and PPD acknowledge assistance from Dr. S. Sivaram, National Chemical Laboratory, India, Dr. Olav M. Aagaard and Mr. Wil J. Belt, DSM Research, The Netherlands. Thanks are also due to Ms. Rituparna Singha, Bhuwneesh Kumar and Mr. Shambhu Bhattacharya for helping the authors to prepare the manuscript.

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Infrared Spectroscopy of Rubbery Materials

5

Infrared Spectroscopy of Rubbery Materials Prajna P. De

5.1 Introduction Linear amorphous polymers can behave as either Hookian elastic (glassy) materials, or highly elastic (rubbery) substances or as viscous melts according to prevailing temperature and time scale of experiments. The different transitions as shown schematically in Figure 5.1 are manifestations of viscoelastic deformations, which are time dependent [1].

Figure 5.1 Temperature dependence of log (shear) modulus in a polymer system showing molecular mechanism of the deformations taking place at different point Reprinted with permission from P. Ghosh, Polymer Science and Technology of Plastics and Rubbers, Tata McGraw-Hill Publishing Company Limited, New Delhi, 1990, Chapter I. Copyright 1990, McGraw Hill Publishing Company Ltd., New Delhi

167

Spectroscopy of Rubbers and Rubbery Materials The deformation in the transition regions is often mechanically reversible, but owing to time dependency, thermodynamically irreversible. It is apparent from the diagram that the useful plastic properties may be exhibited by polymers of wide molecular weight range, while useful rubbery properties are exhibited by polymers which are essentially highly polymeric in nature. To exhibit the rubberiness, the ambient conditions must be far above the glass-to-rubber transition temperature (Tg) of polymer. Tg for elastomers (or rubbers) occurs in the range of –20 to –100 °C, depending on the technique and test conditions. This is manifested in polymers of low cohesive energy density and poor molecular symmetry, which results in an amorphous state, at least in the unstrained state, permitting enough freedom of molecular motion so that the deformation of high magnitude takes place rapidly. Rubber has the ability to extend its length, several hundred percent on application of stress, with virtually immediate and complete recovery on release of the stress [2]. Rubbery behaviour is shown by many polymeric materials, and the first to be exploited was natural rubber (NR). Discovery of Ziegler-Natta catalyst systems, led to the availability of a new generation of synthetic rubbers with controlled molecular architecture. In subsequent years the rapid growth and availability of thermoplastics and in particular polyethylene and plasticised polyvinyl chloride (PVC), led to the replacement of many rubbers in many such applications [2]. In the 1960s and 1970s some conventional rubbers were replaced by thermoplastic rubbers for a number of applications where flexibility rather than high elasticity was the primary concern. The common feature of the thermoplastic rubbers is that the chains are held together at normal ambient temperatures by heat fugitive crosslinks or physical bonds which become ineffective at elevated temperatures [3]. Such crosslinks may be obtained in a number of ways of which the most important are ionic crosslinks, hydrogen bonds and formation of block copolymers. The rubberiness of plasticised PVC is due to hydrogen bonding between the polymer chains and ester plasticiser molecules [3], which by spacing the polymer molecules depress the Tg of PVC. Low density polyethylenes (PE) are rubbery to some extent and the rubberiness becomes more evident when ethylene is copolymerised with a small amount of vinyl acetate. Since infrared (IR) spectroscopy is one of the most widely used techniques for the identification of materials at the molecular level, it has been extensively used to characterise the rubbery materials. In this chapter the rubbery materials encompass PE, plasticised PVC, thermoplastic elastomers and ionomers.

5.2 Polyethylenes PE is a polymer having the simplest structure (CH2)n, with a certain amount of chain branching [4]. As shown in Figure 5.2, the polymer shows strong bands at 2927, 1467 and 724 cm-1, assigned for C-H stretching, deformation and rocking modes of –CH2 groups.

168


Figure 5.2 Infrared spectra of low density polyethylene (LDPE) Reprinted from [53]

The significant chain branching due to the presence of methyl groups gives a peak at 1370 cm-1 overlapped by a methylene band at 1365 cm-1. LDPE from high pressure polymerisation [4] contains vinylidene groups, which absorb at 890 cm-1, while polymers prepared by Ziegler type catalysts [4] show three types of unsaturation: a vinyl group at 910 cm-1, -CH=CH- trans at 965 cm-1, and defect structure of the occasional double bond at 1640 cm-1. At room temperature, PE is a semi-crystalline plastomer (a plastic which on stretching shows elongation like an elastomer), but on heating crystallites melt and the polymer passes through an elastomeric phase. Similarly, by hindering the crystallisation of PE (that is, by incorporating new chain elements), amorphous curable rubbery materials like ethylene propylene copolymer (EPM), ethylene propylene diene terpolymer (EPDM), ethylene-vinyl acetate copolymer (EVA), chlorinated polyethylene (CM), and chlorosulphonated polyethylene (CSM) can be prepared. Studies on structural changes and orientation process occurring in PE during a stressstrain experiment show that the changes in PE films during elongation are manifested in the 730/720 cm-1 (CH2 rocking) absorption bands [5]. The transition moment vectors of the corresponding vibrations are parallel to the a and b axes of the crystalline cell. The samples used were blown PE films having an initial orientation in unit cell c and a axes,

169

Spectroscopy of Rubbers and Rubbery Materials which possess a preferred orientation perpendicular to the film plane and parallel to the machine direction (which is identical to the stretching direction). As far as structural changes during the stretching process are concerned, the orthorhombic structure of crystalline PE is partially transformed into a monoclinic structure as observed by the appearance of a characteristic absorption band at 716 cm-1. Faster stretching velocities increase proportion of the monoclinic structure [6]. It is observed that the bands at 720, 731 cm-1, are resolved to 716, 731 cm-1 with faster stretching. The bands at 720 and 731 cm-1 are due to rocking [7] vibration giving a typical orientation. IR spectroscopy is used to find out the dichroic ratio (R), which can be found out by stretching a PE film, in which, direction of stretch is parallel to slit and then is made perpendicular to the stretching direction. This provides two values of absorption, namely A-parallel and A-perpendicular. The dichroic ratio (R) is defined as: R=

IR absorption for parallel radiation IR absorption for perpendicular radiation

= A|| A⊥

The R values for PE at 731 and 720 cm-1 are 8.9 and 10.0, respectively. Dynamic two-dimensional Fourier transform-IR (FT-IR) spectroscopy has been used to study the nature of the interphase in LDPE [8]. A two-dimensional correlation analysis of the dynamic spectra indicates that neat LDPE is comprised of three regions: an ordered crystalline region, a disordered liquid-like region and a crystal/amorphous interfacial region. The peak which occurs at ~1455 cm-1 has been assigned to methylene bending vibrations in conformationally disordered groups [9], while the peak at ~1467 cm-1 is assigned to bending vibrations in all trans methylene chain sequences outside the crystal structures [10]. Figure 5.3a shows the –CH2 bending and rocking regions of the conventional absorbance spectrum of neat LDPE, as received, whereas Figure 5.3b shows the in-phase and quadrature dynamic step-scan spectra in the CH2 bending region. The dynamic spectra contain only the changes in the absorption due to the application of the mechanical perturbation. The in-phase spectrum in the CH2 bending region is similar to the normal absorbance spectrum [8]. The in-phase and quadrature spectra (a 2-dimensional spectra, which contains just as much area of a certain square) can be used to generate synchronous (Figure 5.3c) and asynchronous (Figure 5.3d) correlation plots [8]. The synchronous plot shows strong auto-peaks at 1464 and 1472 cm-1, which typify crystalline species in semicrystalline polymers. The composite nature of the CH2 bending region is well demonstrated by the asynchronous plot, which shows that 1455 to 1475 cm-1 region splits into at least four independent components at 1458, 1462, 1467 and 1472 cm-1. Analysis of crosspeaks shows that the two bands at 1462 and 1472 cm-1 are out-ofphase with bands at 1458 and 1467 cm-1. This conclusion is in agreement with an earlier work where the 1458 cm-1 band is assigned to the amorphous regions comprising approximately 40% amorphous material, found in the core of the film; whereas skin is

170


a)

c)

b)

d)

Figure 5.3 (a) Static IR absorbance spectrum of unoriented neat LDPE, (b) In-phase and quadrature step-scan dynamic spectra of neat LDPE in CH2 bending region, (c) 2dimensional FT-IR synchronous correlation plot of neat LDPE in the CH2 bending, (d) 2-dimensional FT-IR asynchronous correlation plot of neat LDPE in CH2 bending region Reprinted with permission from A. Singhal and L.J. Fina, Polymer, 1996, 37, 12, 2335. Copyright 1996, Elsevier Science Ltd

171

Spectroscopy of Rubbers and Rubbery Materials crystalline [9] and the 1467 cm-1 band is assigned to methylene bending modes organised in fully transplaner segments in the interphase region [10]. The addition of talc and ethylene vinyl acetate (EVA) to LDPE does not change the dynamic deformation mechanism of the LDPE but their morphological distribution in LDPE is found to be very different [8]. Talc associates intimately with the crystallites and EVA associates with all three of the morphological phases (crystalline, interphase and amorphous). Khastgir and co-workers studied interactions between EVA copolymer and PE with the help of IR spectroscopy [11, 12]. The spectrum of neat EVA shows two distinct carbonyl peaks at 1746 cm-1 and 1729 cm-1. The 1746 cm-1 peak is assigned to the >CO stretching vibration of vinyl acetate group and the peak at 1729 cm-1 is assigned to the keto carbonyl stretching vibration. The formation of ketone [11] group in neat EVA during processing is schematically represented by the reaction Scheme 5.1.

Scheme 5.1 Formulation of ketone in pure EVA during processing Reprinted with permission from I. Ray, S. Roy and D. Khastgir, Polymer Bulletin, 1993, 30, 685. Copyright 1993, Springer-Verlag

But it is interesting to note that the 50/50 EVA/LDPE blend processed under the same conditions shows a single well-resolved >C=O stretching band at 1740 cm-1 (Figure 5.4) which indicates that there is no trace of keto carbonyl groups. The absence of the keto group in the blend may be ascribed to the reaction of the macroradical (formed during processing through abstraction of α-H atoms of EVA) with vinylidene group of LDPE [13]. Hence it may be concluded from IR analysis (Figure 5.4) that under the action of heat and mechanical shearing at the processing condition, LDPEg-EVA is formed [11] according to Scheme 5.2. The grafting reaction takes place generally in the amorphous region of polymers [11]. Similarly, other polar groups like dibutyl maleate (DBM) [14, 15], vinyl trimethoxysilane (VTMO) [16, 17], vinyl triethoxy silane (VTEO) [16, 17, 18] can be grafted onto PE, to make cable sheathing compounds. Sen and co-workers [15] characterised low halogen and non-halogen fire-resistant low smoke (FRLS) cable sheathing compounds from blends 172


Figure 5.4 IR spectra of EVA (———) LDPE (-o-o-) 50/50 EVA/LDPE blend (……) processed at 170 °C, LDPE processed at 130 °C (-..-..-) Reprinted with permission from I. Ray, S. Roy and D. Khastgir, Polymer Bulletin, 1993, 30, 685. Copyright 1993, Springer-Verlag

Scheme 5.2 Reaction of EVA with vinylidene group of LDPE Reprinted with permission from I. Ray, S. Roy and D. Khastgir, Polymer Bulletin, 1993, 30, 685. Copyright 1993, Springer-Verlag

173

Spectroscopy of Rubbers and Rubbery Materials of functionalised polyolefins and PVC by IR spectroscopy. On grafting DBM onto PE, the peak at 1738 cm-1 due to the carbonyl group of ester interacts with PVC through hydrogen bonding, which is reflected by its shifting to lower wavelength [15]. Dipoledipole interactions of the type >C=0 .…. Cl-C- may also exist [15], as shown in Scheme 5.3.

Scheme 5.3 Mechanism of compatibilisation between PE grafted onto DBM (PEgDBM) and PVC through hydrogen bonding and dipole-dipole interaction Reprinted from [15] with permission of John Wiley and Sons, Inc, 1991

Vinyl silanes are grafted onto PE for the modification of properties such as adhesion and dyeability, as well as for preparing compatibiliser for multicomponent polymer blends [16, 17]. The VTMO and VTEO graft co-polymers of PE were prepared by reactive processing in a Brabender extruder in the temperature range of 150-200 °C in the presence of dicumyl peroxide (DCP) [16]. IR spectroscopy has been used to characterise the grafting (Figure 5.5) showing new peaks in the range 800-1200 cm–1, such as 1190 cm-1 (-CH3 rocking in Si-O-CH3), 1094 cm-1 (the anti-symmetric Si-O-C stretching) and 800 cm-1 (the symmetric Si-O-C stretching). Moisture catalysed crosslinking of the silane-grafted copolymer is believed to follow Scheme 5.4. The appearance of a shoulder at 1025 cm-1 due to Si-O-Si stretching indicates crosslinking [16]. Research on the application of radiation processing of polymers, particularly in the food industry necessitates studies on the effects of radiation on various polymer films. This is because different types of radiation may cause different structural changes, rearrangement and fission or formation of chemical bonds in the polymeric film. As a result, the physical,

174


Scheme 5.4 Mechanism of hydrolysis and crosslink formation through the condensation reaction of silanol groups Reprinted with permission from A.K. Sen, B. Mukherjee, A.S. Bhattacharyya, P.P. De and A.K. Bhowmick, Journal of Applied Polymer Science, 1992, 44, 1153. Copyright 1992, John Wiley and Sons, Inc.

chemical, microbiological and mechanical properties of the polymer may suffer some dramatic changes. Chaki and co-workers used electron beam irradiation crosslinking of PE film for the cable sheathing materials [18]. PE (100 parts by weight) was mixed with 5 parts by weight of methyl methacrylate (MMA) at 120 °C and subsequently exposed to electron radiation of different doses (up to 20 Mrad) to prepare PE/MMA graft copolymers. Evidence of successful grafting was verified by IR spectroscopy. A sharp peak around 1030 cm-1 is ascribed to the formation of the methacrylate group in the irradiated crosslinked sample. The grafting level was determined from the absorbance ratio of the peak at 1030 cm-1 to

175


Figure 5.5 IR spectra (a) PE, (b) PE-g-VTMO Reprinted with permission from A.K. Sen, B. Mukherjee, A.S. Bhattacharyya, P.P. De and A.K. Bhowmick, Journal of Applied Polymer Science, 1992, 44, 1153. Copyright 1992, John Wiley and Sons, Inc.

that at 1450 cm-1 (due to –CH2 stretching). The methacrylate concentration of the irradiated samples increased linearly with increase of MMA concentration. Similarly Abou Zeid and co-workers made a detailed study on LDPE film, irradiated with different types of radiation such as 60Co γ-rays, thermal and fast neutrons, and electron beam irradiation [19]. The structural changes of PE films with thickness varying from 8 to 24 μm, were characterised by FT-IR spectroscopy. The characteristic FT-IR absorption bands of unirradiated, γ-irradiated and reactor irradiated LDPE film are presented in Table 5.1. The band at 1018 cm-1 is due to the presence of occulted oxygen in PE film (thickness, 22-24 μm). The bands at 452 and 456 cm-1 appear only for LDPE (thickness, 22-24 μm) which are attributed to the β (C≡C-H). The disappearance of a small band at 888 cm-1 is possibly due to >CH2 out of plane deformation. This may be explained on the basis of evolution of hydrogen gas as a result of the degradation. The intensity of IR absorption bands at 1017, 1113, 1131, 1175, 1241 cm-1 is shown to increase with increasing irradiation dose. The intensity of the bands 1069 cm-1 (C-C stretching), 1287 cm-1 (OH bending), 1377 cm-1 (symmetrical deformation of CH3) (which are initially present in IR spectra of unirradiated films) increase with absorbed dose of radiation, whereas the intensity of the sharp band at 1456 cm-1 is not affected at all. The small band at 965 cm-1 is attributed to C-H out of plane deformation of trans-vinylene groups (trans RCH=CHR).

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Table 5.1 The Characteristic FTIR absorption bands of uniradiated, γirradiated and reactor irradiated LDPE film Original Position of FTIR band, cm-1 Bands initially present in unirradiated PE

Bands induced due to γ radiation

Bands induced due to reactor radiation

Observed Shifts, cm-1

Literature

Intensity change on γ radiation

Low γ dose

452

Constant

456

Constant

888

-

Disappeared

Middle γ dose

High γ dose

DisDisappeared appeared

965

967

965(+)

+

965

965

963.98

1018

1017

1018

1068(+)

+

1017

1018

1018

1069

1113

+

1131

1131(+)

1175

1114

+

1131

1131

1130

+

1177

1177

1178

1241

1241

1287

+

1304

1306(+)

+

1304

1302

1288

1377

1378(+)

+

1378

1378

1377

1410(+)

+

1411

1411

1411 1456 1715

1714

3420

3420

1716(+)

Constant

1456

1457

1457

+

1714

1714

1701

+

+: increase Reproduced from [19] with permission of John Wiley and Sons, 2000

The band at 1303 cm-1 is reported to be one of the bands characterising the amorphous phase in PE. It follows that crosslinking causes a decrease in crystallinity. The band at 1411 cm-1 which is evidence for the occurrence of crosslinking due to the methylene

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Spectroscopy of Rubbers and Rubbery Materials deformation influenced by an adjacent carbonyl group, shows a linear relationship over a wide interval of γ-radiation absorbed doses. Similarly, the intensity of the band at 1716 cm-1, which is attributed to stretching vibration of ketonic carbonyl group, increases remarkably with γ-radiation absorbed doses showing two stages of linear dependence. The intensity of the band in the range (3370-3420 cm-1) increases with the absorbed dose showing linear behaviour for LDPE of various thickness. Hence it can be concluded that the influence of γ-radiation on the structure of PE shows degradation, crosslinking and changes in both crystalline and amorphous regions. Radiation induced graft copolymers with polyfunctionally substituted heterocyclic ring derivatives comprise a very interesting class of polymers because of their significant biological and pharmaceutical activity [20, 21]. The use of N-vinyl pyrrolidone (NVP) as hydrophilic grafting monomer is well known in the radiation grafting on LDPE [22], and the grafted polymer LDPE-g-PNVP is used as the starting material for the preparation of biocompatible polymer surfaces. Here grafting of LDPE was done by 60Co gamma source at a dose rate 1.60 Gy/s at 50 °C in a vacuum oven and the grafted polymer was characterised by IR spectroscopy. FT-IR spectroscopy has been used in the investigation of welding crosslinked polyethylene pipes [23]. Three types of crosslinking systems were used namely, peroxide (PEXa), silane (PEXb) and electron beam (PEXc). Scholten and co-workers [23] observed that only PEXa pipes have a satisfactory electrofusion quality. The strength of electrofusion welds of PEXb and PEXc pipes is not acceptable. The most likely explanation for the differences in weld quality is related to the adhesion theory and more specifically to differences in composition. Figure 5.6 shows the infrared spectra of medium density polyethylene (MDPE), PEXa, PEXb and PEXc. While PEXa exhibits an IR spectrum very similar to uncrosslinked PE, PEXb contains a rather high concentration of silicon containing groups, as is obvious from the broad band due to Si-O stretching and bending modes around 1080 cm-1. These groups are residues of the crosslinking process. The large band at ~1740 cm-1 indicates that PEXc contains EVA, which is supported by the occurrence of smaller bands at other positions (not shown). These polar ‘impurities’ may have a deteriorating effect on the quality of electrofusion welds produced with PEXb and PEXc. In PE/aluminium laminates, which are used for packaging, the bond strength can be improved by surface oxidation of the polymer. An alternative to surface oxidation is to use copolymers of ethylene and monomers containing polar groups that can provide stronger bonds to metal thus increasing adhesion [24, 25]. The effects of three functional groups in ethylene copolymers, namely EVS [polyethylene-co-vinyl trimethoxy silane], EBA [polyethylene-co-butyl acrylate], EAA [polyethylene-co-acrylic acid) on the adhesion was studied [25]. The interface in polymer/metal laminates has been analysed by FT-IR

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Figure 5.6 Part of IR spectra of PEXa, PEXb and PEXc are compared with that of MDPE [bending vibrations at ~1450 cm-1 omitted] Reprinted from [23] with permission of Elsevier Applied Science, on behalf of the Institute of Materials, New York, 1998

spectroscopy. It is observed that the increased amount of free acid at the interface can be attributed to interfacial hydrogen bonds between the hydroxyl groups of the acid and the surface AlOH groups, forming Al-O-Al groups. Furthermore, the band at 1738 cm-1 represents the carbonyl group in an interfacial ester link between the polymer and the metal originating from a reaction between surface AlOH groups and hydroxyls of acid groups. The carbonyl of ester copolymer gives two absorptions in the IR spectrum. The downshift of the carbonyl frequency at the interface can be attributed to interaction with the aluminium oxide of the substrate. Ishida and co-workers reports that in the case of EVS, the silane-related bands in the transmission spectrum are the same as that in the case of silane modified polymer [25]. Examples are the bands at 1190 cm-1 (for –CH3 rocking in Si-O-CH3), 1094 cm-1 (anti-symmetric stretching of Si-O-C) and 800 cm-1 (the symmetric stretching of Si-O-C). In addition the appearance of a shoulder at 1025 cm-1 representing the anti-symmetric Si-O-Si stretching, indicates formation of silanol groups followed by condensation to siloxanes. Due to crosslinking via the silane groups, the adhesive strength for EVS is higher than that of EBA and EAA.

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Spectroscopy of Rubbers and Rubbery Materials The production of polymers with enhanced physical properties such as high modulus fibres has led to detailed investigations on the polymer morphologies. In several cases PE single crystals have been used as well characterised model starting materials for studies of deformation. Recently the isotopic labelling technique has been used to establish the conformation of individual polymer molecules in solution-grown single crystals of PE [26]. Okoroafor and co-workers have prepared mixed crystals [26] of normal polyethylene (PEH) and fully deuterated polyethylene (PED) from dilute solution of mixed xylenes at 70 °C and characterised them by using the mixed crystal IR technique. They have compared the CX2 bending vibration (X=H or D) for PEH and PED mixed crystal of PE and observed reduction in the absorbance of the outer CD2 bending components, which provides evidence of lamellar break up. Characterisation of ethylene in ethylene-propylene copolymers (EP) and in isotactic polypropylene (iPP)/EP systems by deconvolution of IR spectra was studied by Costa and coworkers [27]. The method is based on the deconvolution of the spectrum in the CH2-rocking range, i.e., 800-680 cm-1, where the bands of interest overlap. The overlapping bands at 720 and 730 cm-1 are due to ethylene crystallinity. Similarly a band for amorphous PE appears at 722 cm-1, a band for single ethylene group between two propylene groups, appears at 735 cm-1, a band at 752 cm-1 is assigned to –(CH2)2- between tertiary carbon atoms, for head-to-head and a band at 770 cm-1 is assigned to pendent ethyl groups. The six bands present in this region are signal averaged in position and width (Figure 5.7). The spectra are then deconvoluted (curve fitted, Figure 5.7) assuming a Lorentzian shape for bands. The band at 1167 cm-1 (with a shoulder at 1156 cm-1) corresponding to a CH3 vibration is considered as an internal standard. The method is checked by varying some fitting parameters. The amounts of total (-CH2-CH2) units, isolated ethylene and structural defaults in PP were determined for different iPP/EP blends.

5.3 Polyvinyl Chloride Like other polymers, the structure of PVC is identified by IR spectroscopy [4]. Figure 5.8 shows the spectra of PVC, indicating the presence of saturated C-H groups with CH2 deformation band at 1430 cm-1, which has moved about 30 cm-1, suggesting the presence of polar substituents. The C-Cl stretching bands are observed in the range of 600-700 cm-1. Usually two different polymers do not mix at the segmental level but a favourable interaction between the two polymers can allow one to obtain homogeneous blends. FT-IR is a potential tool for the investigation of the mutual compatibility of various polymers. The small spectral changes due to these interactions can be detected by this method. If two polymers are immiscible, one can synthesise a spectrum of the blend by co-adding, in the appropriate

180


Figure 5.7 Deconvolution of an iPP/EP sample in the range 800-680 cm-1 Reprinted from [27] with permission of John Wiley and Sons, Inc., 2000

Figure 5.8 Infrared spectra of PVC Reprinted with permission from W.F. Maddams in Analysis of Polymer Systems, Eds., L.S. Bark and N.S. Allen, Applied Science Publishers, London, 1982, Chapter 3, 59. Copyright 1982, Bark and Allen

181

Spectroscopy of Rubbers and Rubbery Materials proportions, the spectra of the two pure components, which is then compared with the observed spectrum of the blend. On the other hand, when polymers are miscible, only one phase exists and there should be spectral differences between the co-added spectra of the pure components [5]. Compatible and incompatible polyester – PVC blends have been considered. Examples include systems, poly-∑-caprolactone (PCL)/PVC which are compatible [5, 28], in the melt and exhibit partial compatibility in solid state and poly-β-propiolactone (PPL)/PVC blends which are known to be incompatible [5, 28]. Figure 5.9 shows the infrared spectra of the carbonyl stretching vibration (in the range 1600-1800 cm-1) for the different blends.

Figure 5.9 (a) Spectra of PVC/PCL blends at room temperature for (A) neat PCL (B) 1:1 (C) 2:1 (D) 3:1 (E) 5:1 (F) 10:1, (b) Spectra of PPL and PPL/PVC (30/70) blend at room temperature Reprinted with permission from Developments in Polymer Characterisation, Ed., J.V. Dawkins, Applied Science Publishers, London, 1983, Chapter 3, 119. Copyright J.V. Dawkins, 1983

182

Infrared Spectroscopy of Rubbery Materials Specific interactions between PCL and PVC are clearly indicated. In the solid state (Figure 5.9a) the spectrum of neat PCL indicates the presence of crystalline (1724 cm-1) and amorphous (1737 cm-1) bands. At mole ratios up to 2:1 of PVC to PCL, the spectra indicate that in the solid state the blends consist of crystalline and amorphous phases. As the PVC concentration increases, a parallel increase of the intensity of the amorphous band is observed. Moreover, the frequency shifts observed for both the crystalline and amorphous bands as a function of the composition of the blend suggests that specific interactions between the two polymers occur. No shift is observed in the carbonyl stretching vibration of PPL/PVC blends, in the molten state or in the solid state over the entire range of compositions and the two polymers are incompatible [28]. PVC and acrylonitrile butadiene rubber (NBR) blends undergo self-crosslinking [29-31] during high temperature mixing in the absence of any external curing agents. Evidence of the chemical interaction is confirmed by IR spectroscopic studies, which show that the reaction takes place through the amide and acid groups formed by the hydrolysis of the -C≡N groups, in the presence of hydrogen chloride liberated from PVC [32]. The interaction [31] of PVC with nitrile rubber is shown in Scheme 5.5.

Scheme 5.5 Chemical Interactions of PVC and nitrile rubber during mixing Reprinted with permission from N.R. Manoj and P.P. De, Polymer, 1998, 39, 3, 733. Copyright 1998, Elsevier Science Ltd

183

Spectroscopy of Rubbers and Rubbery Materials The representative IR spectra for PVC and NBR blends [31] are shown in Figures 5.10 and 5.11 [31]. Neat PVC shows two sharp absorbance peaks at 1734 and 1590 cm-1 which are due to the carbonyl groups formed during the manufacturing stage [33] and the polymer backbone defect [34].

Figure 5.10 IR spectra of PVC and NBR before mixing (a) PVC (c) NBR after mixing for 60 minutes at 180 °C, (b) PVC (d) NBR Reprinted with permission from N.R. Manoj and P.P. De, Polymer, 1998, 39, 3, 733. Copyright 1998, Elsevier Science Ltd

184


Figure 5.11 Infrared spectra of 50/50 PVC/NBR blends after mixing (a) for 5 minutes, (b) for 60 minutes (stabilised with tribasic lead sulfate) at 180 °C, (c) at 160 °C, (d) at 180 °C, (e) at 200 °C for 60 minutes Reprinted with permission from N.R. Manoj and P.P. De, Polymer, 1998, 39, 3, 733. Copyright 1998, Elsevier Science Ltd

During processing, PVC undergoes thermooxidative degradation leading to the formation of conjugated polymers, which absorb at 1715-1660 cm-1 The broad peak at 3280 cm-1 may be due to the formation of –OH groups during degradation [31]. The difference in the

185

Spectroscopy of Rubbers and Rubbery Materials spectra of NBR before and after processing (Figure 5.10) is the slight broadening of the band at 3510 cm-1 due to formation of –OH groups during degradation, and a new peak at 1575 cm-1 due to formation of triazine generated by cyclisation of adjacent nitrile groups [31]. The absorption band at 2224 cm-1 is due to the stretching vibration of C≡N groups. Figure 5.11 shows the spectra of 50/50 PVC/NBR blends at different temperatures. A comparison of the spectra of the blends mixed for 5 and 60 minutes at 180 °C (Figures 5.11, a and d) show the changes occurring during crosslinking. The broadening and the increased absorption at 3500-3200 cm-1 show the formation of O-H and N-H groups in the system due to hydrolysis of the nitrile groups. At 160 °C the peak is smaller and at 180 °C, there is increased absorption indicating the temperature dependence of the reaction. At 200 °C, the peak intensity decreases, indicating that some of the –OH groups react during mixing. In spectra (Figure 5.11b) the blend is mixed with stabiliser, it shows the same three peaks of PVC (without stabiliser) observed with almost equal intensity within 1715-1660 cm-1. The intensity of these peaks change, when they are heated without stabiliser as seen in Figures 5.11c, 5.11d, 5.11e. The peak centred at 1720 cm-1 represents the carbonyl groups of aliphatic acids, esters and ketones. The carbonyl groups in amides absorb at 1630 cm-1 and carbon-carbon double bonds absorb at 1660 cm-1. During crosslinking reaction, the concentration of carbonyl groups increase due to formation of amides, acids and esters Figures 5.11 d and 5.11e and that of carbon-carbon double bonds increase due to the simultaneous degradation (and e). This is manifested in the broadening of the peaks and increased absorption as the processing temperature and time increases. The changes in the concentration of nitrile group are not visible since the absorption is sharp and strong. But in order to confirm any change of the intensity of the band for the acrylonitrile (-C≡N) group, Manoj and co-workers [31] have mixed separately NBR, hydrogenated acrylonitrile butadiene rubber (HNBR) and PVC/NBR (50/50 blend) and PVC/HNBR (50/50) blend at 180 °C for 60 minutes in the Brabender Plasticorder and IR spectra of each compound was taken. The thermal treatment does not produce any perceptible change in the intensity of the band at 2224 cm-1, but the decrease in —CN concentration in the blend is attributed to partial hydrolysis of cyanide groups by the splitting of hydrogen chloride from PVC during heating in the presence of moisture [31] (Scheme 5.5). De and co-workers studied chemical changes in the hot air and fuel ageing of PVC/NBR and PVC/HNBR blends with the help of infrared spectroscopy [35]. The thermal stability of PVC blended with polymers acting as impact modifiers depends mainly on the type of polymer, blend morphology and degradation environment. The effect of the type of epoxidised butadiene-styrene block copolymer [ESBS, linear butadiene/ styrene (B/S) or radial (E(B/S)n) containing 0%-27% of epoxy groups] on thermal dehydrochlorination of PVC-ESBS blends has been investigated by Meissner and coworkers [36] in the temperature range 170-180 °C under non-oxygen atmosphere. Thermal stability of PVC-ESBS blends is estimated on the basis of induction time and maximum rate of hydrochloride emission from the system. On the basis of the analysis

186

Infrared Spectroscopy of Rubbery Materials of the IR spectra of PVC and its mixtures with epoxidised co-polymers before and after heating at 180 °C in the range 400-4000 cm-1, it is concluded that hydrogen chloride liberated from PVC reacts with double bond and epoxy groups to form –CH2-CHCland –CH(OH)-CHCl- as indicated by the bands at 3600 cm-1 (OH) and 698 cm-1 (C-Cl). Similarly IR spectra has been used to find out the thermal stability of PVC-diocytyl phthalate plastigels in presence of zinc stearate (ZnSt2) and epoxidised soyabean oil (ESO) [37]. ZnSt2 added to PVC is characterised by an absorption band at 1540 cm-1 corresponding to the stretching vibration of carboxylate groups. The intensity of carboxyl band is proportional to the concentration of ZnSt2, and it decreases with increasing heating time, as it reacts with hydrogen chloride liberated from PVC. The OH group of organic acid is characterised by an absorption band at 3400 cm-1, corresponding to the stretching vibration. The best heat stability and gelation effects are obtained with ESO for long heating times, which are confirmed by dehydrochlorination rate constants. Like thermal ageing, photoageing of PVC has been studied by Verdu and co-workers [38-41]. It is widely recognised that PVC photoageing is a very complex process due to the co-existence of two (sharply interrelated) chain reactions, namely, oxidation and hydrogen chloride elimination. Photoelimination creates chlorine (Cl) radicals able to initiate oxidation chains, HCl able to catalyse hydroperoxide or peroxy decomposition and polyenes able to play a photosensitising or photostabilising role depending essentially on their length [39, 40]. Verdu and co-workers [39] have used FT-IR spectra to determine the carbonyl concentration at 1720 cm-1 and have observed that a significant decrease of number average molecular weight and the highest rate of carbonyl build up take place in the superficial layer of about 50 μm thickness, whereas the subcutaneous layer of about 200-300 μm thickness is characterised by a predominant crosslinking and conjugated polyene growth, when it becomes insoluble and highly discoloured. The accelerated photoageing of non-photostabilised, non-pigmented bulk PVC samples have been studied in photochemical reactors at different temperatures (40-70 °C) with sources of different intensities starting from 7.2 x 1018 photons per cm2h-1. As was observed in earlier studies [36, 37], they have observed the existence of two layers [40]: the superficial layer C = C in NR (Figure 5.12) is shifted to 836 cm-1. Similarly the peak at 1370 cm-1 (C-H stretching of CH3 group) shifts to 1373 cm-1, while the peak for C=C double bond shifts from 1660 cm-1 to 1658 cm-1, and the band at 1467 cm-1 for –CH2 in LDPE (Figure 5.12) is shifted to 1462 cm-1. The spectra thus confirm that there exist only physical interactions in NR-PE blend. Polyurethane TPE are block copolymers consisting of alternating hard and soft segments, separated into two phases [54] due to thermodynamic incompatibility.

188


Figure 5.12 Infrared spectra (a) LDPE, (b) NR, (c) NR/LDPE (50/50 blends) Reprinted from [53]

Two series of polyether polyurethanes (PU) based on hydroquinone bis (β-hydroxyethyl) ether (HQEE) or 1,4-butanediol (BDO) as a chain extender were prepared by the one step bulk polymerisation process. By varying the mole ratio of poly tetra methylene oxide (PTMO) extender (with Mn = 1000 and Mn = 2000) and 4,4´-diphenylene methane diisocyanate (MDI) the two series of HQEE (PU1000H1, PU 1000H2, PU2000H1,

189

Spectroscopy of Rubbers and Rubbery Materials PU2000H2, PU2000H3) and BDO (PU1000B1, PU1000B 2, PU2000B1, PU2000B2, PU2000B3) were prepared [55]. Evidence of hydrogen bond formation is obtained by means of FT-IR spectroscopy [55]. The hard segment units derived from MDI react with (a) HQEE and (b) BDO according to Scheme 5.6.

Scheme 5.6 Structure of HQEE-MDI and BDO-MDI Reprinted with permission from L. Zha, M. Wu and J. Yang, Journal of Applied Polymer Science, 1999, 73, 2895. Copyright 1999, John Wiley and Sons, Inc.

In both cases, the N-H group of the urethane serves as the proton donor, while the acceptor includes the carbonyl oxygen in the urethane group. It is generally agreed that C=O stretching region is favourable for the estimation of hydrogen bonding in PU, qualitatively and quantitatively. From the spectra, it can be observed that the absorption band at 3451 cm-1 associated with the free NH groups, seems to be absent in HQEEbased PU. This indicates that approximately 100% N-H groups in HQEE-based PU are hydrogen bonded. The carbonyl peak of urethanes splits into the following two peaks: one at about 1702 cm-1 results from self-association of NH with C=O in the interior of the hard domains; and the other peak at 1732 cm-1 is related to the free carbonyl groups present in the mixed soft segment phase. To assist in probing the complex by hydrogen bonding structure from the IR spectra of the PU, the technique of Fourier deconvolution [55] has been utilised. The deconvoluted spectra for carbonyl stretching region are shown in Figure 5.13. As can be seen from Figure 5.13 the lowest frequency associated with ordered hydrogen bonded carbonyl in crystallite for HQEE based PU is lower by 4 cm-1 than BDO-based PU (1698 and 1702 cm-1) [56]. This indicates that the former shows a higher extent of ordering and a stronger ordered hydrogen bond. Compared with the BDO-based hard segment, the presence of benzene ring (-O-Ph-O-) in the HQEE based hard segment

190


Figure 5.13 Deconvoluted FTIR spectra in the C=O stretching region of (a) HQEE based PU and (b) BDO based polyurethanes Reprinted from [55] with permission of John Wiley and Sons, Inc., 1999

191

Spectroscopy of Rubbers and Rubbery Materials contributes to the ordering due to chain interaction. Figure 5.13 shows that two types of disordered hydrogen bonded carbonyls are observed at frequencies intermediate to the free and ordered hydrogen-bonded carbonyls for HQEE-based PU, whereas only one is observed for BDO-based PU. As were assigned previously [56], for HQEE based PU and BDO based PU the band at around 1715 cm-1 is associated with the hydrogen bond that is short-range in the amorphous phase of the hard domains or the interfacial phase region, whereas the band at about 1708 cm-1 corresponds to the long-range hydrogen bond in the amorphous phase of the hard domains. With an increase in the hard segment content for HQEEbased polyurethanes the absorption band at 1732 cm-1 associated with free carbonyl diminishes, while the band at 1698 cm-1 increases dramatically among hydrogen-bonded carbonyl stretching bands. Qualitatively, the spectra for hydrogen bonded carbonyl is dominated by the band at 1708 cm-1 at a low hard segment content. As the hard segment content is raised, the dominant band is replaced by 1698 cm-1 components. Hence it can be concluded that compared with the BDO-based hard segment, the (-O-Ph-O-) linkage in HQEE makes the hard segment more incompatible with polyether soft segment and the cohesion in the hard domain enhances. Hydrogen bonding plays a critical role in determining the morphology and properties of polyamides and PU and polyurethane-ureas. Yilgor and co-workers studied the hydrogen bonding in polydimethyl siloxane and polyether based urethane and urea type segmented copolymers with IR spectroscopy [57]. The extent of hydrogen bonding in PU or polyureas can be qualitatively studied by determining the frequency shifts in a hydrogen bonded (-NH-) and (-C=O) peaks, that is the peak due to (-N-H…. O=C) relative to the free (-N-H) and (-C=O) peaks. Table 5.2 gives a detailed list of absorption frequencies for various groups and their hydrogen bonded complexes. Comparative FTIR spectra for (-N-H) and (-C=O) stretching regions of model urethane and urea compounds and the copolymers are given in reference [57]. Similarly hydrogen bonding in crosslinked urethane/urea polymers with two soft segments, polypropylene oxide and polybutadiene, have been studied by Zhao [58] and co-workers with the help of IR spectra. Furthermore, IR analysis has been used for characterisation of a series of segmented block copolymers of NR and 1,3 butene diol-toluene diisocyanate [59] and bisphenol A-toluene diisocyanate oligomers [60]. Like PU TPE, blends of thermoplastic polyurethanes and polyamide-12 (PA-12) have been studied by Polosmak and co-workers [61]. They have mixed two types of thermoplastic polyurethane (TPU) based on oligoether (polytetramethylene oxide, molecular weight, 1000) and oligoester (polyethylene butylene glycol adipate, molecular weight, 2000) and PA 12 were characterised by IR spectra and thermal analysis. IR spectra of TPU, PA-12 and their blends show that in amide one (A1) carbonyl absorbancy is seen to split [55] into two main bands with maxima at 1705 and 1730 cm-1. At 1730 cm-1,

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Table 5.2 Characteristic IR absorption frequencies for polyurethanes and polyureas Group

Mode

Frequency (cm-1)

N-H

free

3445-345 0

N-H

N-H … N-H

3315-334 0

N-H

N-H … E (ether)

3260-329 0

C-O (urethane)

free

1730-174 0

C=O (urethane)

C=O .. H-N

1730-1710

C=O (urea)

free

1690-170 0

C=O (urea)

C=O… H-N

1660-1670 (disordered)

C=O

C=O…H-N

1630-1645 (ordered )

NH-C=O

Amide II

1540-1560

Reprinted from [57] with permission of Elsevier Science Ltd., London, 1999

absorption bands of free carbonyl groups of the urethane fragments are found along with low energy hydrogen bands, e.g., dimer, trimers, and also hydrogen bonds developing between NH group and oxygen in the oligoester block. The ratio of optical densities of the bands at 1705 and 1730 cm-1 can also show the degree of self-association of urethane fragments. Addition of PA-12 results in variation of the shoulder intensity ratio in A1 band and the extent of this influence depends on the nature of TPU, plastomer and concentration of the latter. In ‘optical blends’ which are obtained by putting PA-12 and TPU films on a surface of germanium crystal, the D1705/D1730 ratio (= optical density of the bands at 1705 cm-1/optical density of the bands at 1730 cm-1) is constant being independent on composition. Due to this, its variation in real blends results from the influence of plastomer on hydrogen interactions in TPU. Addition of PA-12 lowers the D1705/D1730 ratio for ester containing TPU2 (ester TPU), whereas for oligoether containing TPU1 (ether TPU) it increases compared with the virgin thermoplastic elastomers. So in the blends with PA-12, the TPU2 domain structure undergoes partial degradation whereas for TPU1 the extent of micro segregation increases. For two TPU in blends with PA-12, NH is usually shifted to lower frequency region, indicating formation of higher energy hydrogen bonds than that in virgin components (PA-12 and TPU). It is observed that the effect of thermoplastic polymer on hydrogen bonding in TPU depends not only on its ability to form hydrogen bonds with PA-12 macromolecules, but also on the distribution pattern of PA-12 in individual blocks [62]. Jha and co-workers [63, 64] prepared thermoplastic

193

Spectroscopy of Rubbers and Rubbery Materials elastomeric blends of Nylon 6 and acrylate rubber (ACM) and have studied the interaction between the two components with IR spectroscopy [63], which shows reduction in the intensities of the peaks corresponding to epoxy groups of ACM as well as carboxylic acid and amine end groups of Nylon 6, suggesting that a chemical reaction between the above groups takes place at the processing condition. IR spectroscopy has been used to detect the chemical changes during the thermal degradation of PU-epoxide resin (ER) foam [65] and photochemical degradation of aliphatic polyester urethane [66]. According to Zhang and co-workers [65], the thermal stability of the rigid neat polyurethane foam (PUF) is slightly improved when the ER is incorporated. With increase of temperature, two stages of weight-loss process are observed for the rigid PU/ER interpenetrating network foam (IPNF) and for the neat PUF. FT-IR analysis [66] suggests that the first stage of weight loss is due to the degradation of the polyol-derived blocks of the PU and the second weight loss stage is governed by the degradation of the MDI-derived blocks of the PU, along with the degradation of the ER component. FT-IR analysis shows that the ether bond (at 1072cm-1) disappears first at around 300 °C and the urethane group (at 3323 and 1722 cm-1) is finally degraded at about 400 °C. Photooxidation of polyesterurethane takes place through a free radical mechanism [66]. Conventional PU based on aromatic diisocyanates such as MDI and toluene diisocyanate (TDI) are known to undergo yellowing on exposure to UV radiation. It is generally agreed that the principal structural species responsible for discolouration of MDI-based PU are mono and diquinone imide [66]. It is shown by FT-IR spectra that long wavelength irradiations provoke an induced oxidation of urethane functions. This reaction is initiated by hydrogen atom abstraction on the methylene groups in the α-position of nitrogen atoms. On irradiation at short wavelength the results provide evidence for a dual mechanism of photooxidation and induced oxidation, which are characterised by an appreciable loss of ester structure, as indicated by the decline of the bands at 1731, 1263, 1167 and 1065 cm-1. Actually photo scissions involve the homolysis of C-N and C-O bonds producing –COOH- and H2N-COO- groups, and with further irradiation, decarboxylation takes place. Introduction of ionic functional groups, which are pendent on the backbone chain, results in formation of ionomers. In the case of polyblends, intermolecular ionic interactions facilitate compatibilisation [67, 68]. The mechanical properties of the polyblends are dependent on the specific interactions between the component polymers promoting interfacial adhesion [49, 69]. There are several reports on compatible blends, wherein compatibility is induced by ion-ion and ion-dipole interactions [70, 71]. Recently De and co-workers have reported studies on thermoplastic elastomers based on ionomeric polyblends [72-81]. They have characterised the blends with the help of infrared spectroscopy. For example, an ionic thermoplastic elastomer (ITPE) [75] is prepared by melt blending zinc salts of carboxylated nitrile rubber (Zn-XNBR) and maleated high density polyethylene (Zn-mHDPE) according to formulations given in Table 5.3. The synergism in physical properties of the resulting

194


Table 5.3 Formulation of the mixes Ingredient

ph r

Mix number

M0

M1

M2

M3

M4

M5

M6

XNBR

100

90

80

70

60

50

0

m-HDPE

0

10

20

30

40

50

10 0

ZnO

12

12

12

12

12

12

12

Stearic acid

1

1

1

1

1

1

1

Reprinted from [75] with permission of John Wiley and Sons, Inc., 1998 phr: parts per hundred rubber

ITPE is due to the formation of strong intermolecular ionic crosslinks, which act as a compatibiliser. The IR spectra [75] of Zn-mHDPE and Zn-XNBR in the range of 17501250 cm-1 are shown in Figure 5.14. The absence of the band corresponding to hydrogen bonded carboxylic acid pairs (1700-1720 cm-1) indicate almost complete neutralisation of the acid groups in both polymers. The spectrum of Zn-XNBR shows a weak band at 1665 cm-1, which is ascribed to the –C=C— stretching mode and the asymmetric carboxylate stretching region shows a doublet at 1587 and 1541 cm-1. These bands are strong and intense compared to the spectrum of Zn-mHDPE, which shows a doublet at 1596 and 1552 cm-1. The stronger asymmetric carboxylate stretching band observed for Zn-XNBR is due to the presence of a higher proportion of carboxylate ions in Zn-XNBR than in Zn-mHDPE. The splitting of the asymmetric carboxylate stretching band into a doublet is assigned to the different coordinated structures of the zinc cation. The band, which is strong and intense at 1462 cm-1, is due to a –CH2 bending vibration of Zn-mHDPE. The bands observed at 1445 and 1415 cm-1 may also have a contribution from the symmetric carboxylate stretching. A weak band at 1358 cm-1 accounts for –CH2- wagging. IR spectra of ionomeric polyblends (mixes M1, M3 and M5) in the range of 1750-1250 cm-1 (assigned for octahedral [82] structure of zinc carboxylate ion) are shown in Figure 5.14. These spectra show a strong and intense doublet at 1587 and 1541 cm-1 in asymmetric carboxylate stretching region. The band at 1587 cm-1 is assigned to the tetrahedral structure of zinc-carboxylate and the 1541 cm-1 band accounts for octahedral structure. It is interesting to note that as ZnmHDPE content in the blend increases, the intensity of the 1541 cm-1 band decreases as compared to that of 1587 cm-1 band. The intensity of the –C=C- stretching band at 1665 cm-1 decreased with increase in Zn-mHDPE content in the blend. The IR spectrum of M5 shows a clearly resolved band at 1462 cm-1 indicating a –CH2- bending vibration. The difference

195


Figure 5.14 Infrared spectra of (a) Zn-mHDPE, (b) Zn-XNBR, (c) Mix M1, (d) Mix M3, (e) Mix M5 in the range 1750-1250 cm-1 Reprinted from [75] with permission of John Wiley and Sons, Inc., 1998

spectrum [75] in Figure 5.15 is obtained by subtracting the weighted addition spectra of the neat polymers from the observed spectra of the corresponding blends. In the case of incompatible blends, the spectrum of the blend should be similar to that obtained by the summation spectra of the neat polymers [69]. But the compatible blends show marked changes in the spectra. It is evident from the difference spectra that the ionic interaction is stronger in the blends than that in neat polymers. This is supported by the positive absorption band at 1586 cm-1, which is due to the tetrahedral asymmetric COO- stretching band in the difference spectra. It is also noted that the 1552 cm-1 band intensity in the mixes M1

196


Figure 5.15 Difference spectra obtained by substracting the sum of The spectra of the neat polymers from the blend spectra (a) Mix M1, (b) Mix M3, (c) Mix M5 Reprinted with permission from P. Antony and S.K. De, Journal of Applied Polymer Science, 1998, 70, 483. Copyright 1998, John Wiley and Sons, Inc.

197

Spectroscopy of Rubbers and Rubbery Materials and M3 and the 1569 cm-1 band intensity in M5 decrease, indicating a change due to the mutual interaction between the ionic groups present in the ionomers. A series of sulfonated ionomers have been prepared and characterised by IR studies [8385]. Recently Ghosh and co-workers reported the preparation of an ionic thermoplastic elastomer based on the zinc salt of sulfonated maleated EPDM rubber (Zn-s-m-EPDM) [86-89]. They also studied incorporation of semi-reinforcing furnace (SRF) and intermediate super abrasion furnace (ISAF) carbon black and precipitated silica on zinc salt of sulfonated maleated EPDM rubber [88]. The rubber-filler bonding has been studied by IR spectroscopy [89]. IR spectra of zinc stearate, precipitated silica, SRF and ISAF carbon black are shown in Figure 5.16. In Figure 5.16a, zinc stearate shows the characteristic absorbance at 1577 cm-1, which is due to asymmetric stretching of the bridging type of carboxylate groups. The spectrum of precipitated silica is characterised by the strong absorbance at 1096 cm-1, which is due to Si-O-Si asymmetric stretching. The other characteristic peaks at 3760, 960, 798 cm-1 are due to -OH stretching of SiOH, Si-O stretching of Si-OH and Si-OH deformation, while the peak at 470 cm-1 is due to Si-O-Si symmetric stretching [90]. The peaks at 798 and 470 cm-1 for precipitated silica are not shown in Figure 5.16a. SRF carbon black shows characteristic absorbance at 3730 cm-1, which is due to O-H groups in the substituted phenolic compounds [91]. The broad peak at 3162 cm-1 is attributed to composite absorption of hydrogen bonded –OH groups from alcohols, phenols and enols. The peak at 1693 cm-1 is due to carbonyl stretching. ISAF black shows bands at 1350 and 1385 cm-1 due to –OH bending in water and C-O stretching in phenols [92]. The band at 1574 cm-1 confirms the presence of tetra- hydroquinone and the peak around 1634 cm-1 is due to polycyclic quinines [91]. The absorbance bands at 1750 and 3400 cm-1 are due to lactones and hydrogen bonded phenolic –O-H groups. The IR spectra of silica-filled compound is shown in Figure 5.16b, where the peak at 3472 cm-1 is believed to be due to hydrogen-bonded structure in the vicinal (adjacent) groups, whereas the peak at 3666 cm-1 is due to hydrogen bonded structures. The peak in the asymmetric stretching region occurs at 1567, 1548, and 1532 cm-1 and the carboxylate symmetric stretching occurs at 1410 cm-1 The peak at 1189 cm-1 is due to the asymmetric stretching of the sulfonate groups. The split pattern between 1000 and 1200 cm-1 indicates specific interaction of silica filler with sulfonate groups. The ionomeric compositions filled with SRF and ISAF carbon blacks show a broad band in the sulfonate stretching frequency region (800-1400 cm-1). The SRF carbon black shows a broad band at 3225 cm-1 due to a composite absorption by hydrogen bonded O-H groups present in both carbon black and the ionomer [91, 92]. The ISAF carbon black-filled ionomer shows band splitting in the high-frequency region, a carbonyl stretching frequency at 1656 cm-1 and hydrogen-bonded carboxylate anions at 1572 and 1536 cm-1. Similarly, IR spectrum of the zinc stearate filler-loaded ionomers shows peaks at 1585, 1547, 1528 and 1223 cm-1 and thus indicates that an interaction occurs between the active sites of the filler and the zinc stearate-filled ionomer.

198


Figure 5.16 (a) IR spectra of zinc stearate, SRF carbon black, silica, ISAF carbon black, (b) Infrared spectra of Zn-s-m-EPDM containing zinc stearate, SRF carbon black, silica and ISAF carbon black Reprinted from [89] with permission of John Wiley and Sons, 2000

UV curable waterborne coatings made of PU-acrylate ionomers have been synthesised and characterised by IR spectroscopy [93]. Similarly a series of novel PU ionomers with polydioxolane (PDXL) as soft segment is prepared and identified by FT-IR spectroscopy [94].

199


5.5 Conclusion Different types of interaction in rubbery materials, which include PE, plasticised PVC, TPE, ionomeric polyblends, and blends of plastics and rubbers can be characterised by different types of IR spectroscopic techniques.

Acknowledgements The author thanks Professor S.K. De, Professor, Rubber Technology Centre, IIT Kharagpur for his moral support and constant co-operation in proof reading. Thanks are due to Mr. R.S. Rajeev, Mr. Shambhu Bhattacharyya and Mr. Bhuwneesh Kumar for their assistance in preparing the manuscript.

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206

6

Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy Herman G. Dikland and Martin van Duin

6.1 Introduction

6.1.1 General Introduction to EPDM [1-3] EPM is a copolymer, consisting of ethylene and propylene. As a result of the random incorporation of the monomers, EPM with medium ethylene content are amorphous elastomers. EPM with high ethylene contents have diffuse crystalline regions, yielding some rigidity although retaining rubbery characteristics. As a result of the absence of unsaturation, EPM cannot be vulcanised with sulfur and it is not very reactive to peroxide curing. To increase the reactivity for crosslinking a diene, namely 5-ethylidene-2norbornene (ENB) or dicyclopentadiene (DCPD) is terpolymerised in combination with ethylene and propylene. The resulting EPDM combines a saturated polymer backbone with residual unsaturation in the side groups. Typical examples of such polymers are shown in Figure 6.1. EPM and EPDM, generally referred to as EPDM are classically produced via ZieglerNatta catalysis in solution or slurry processes. Over the last decade metallocene catalysis [4, 5] and gas phase technology [6] have been developed. EPDM is produced on a commercial scale in a variety of chemical compositions. Typically, the ethylene content

Figure 6.1 Chemical structure of EPDM with 2-ethylidene-5-norbornene (left) and dicyclopentadiene (right) as diene

207

Spectroscopy of Rubbers and Rubbery Materials ranges from 40 to 80 wt.% and the diene content from 0 to 12 wt.%. The weight average molecular weight of EPDM typically ranges from 100 to 1,000 kg/mole in combination with molecular weight distributions of 2.0 to 10. The degree of long chain branching can be controlled to a considerable extent [7]. As for all elastomers, the processing characteristics of polymers in combination with the compound recipe are used to determine the processing behaviour of the compound and as well as the mechanical, elastic and dynamical properties of the final product after crosslinking. As a result of its saturated polymer backbone, EPDM is more resistant to oxygen, ozone, UV and heat than the low-cost commodity polydiene rubbers, such as natural rubber (NR), polybutadiene rubber (BR) and styrene-butadiene rubber (SBR). Therefore, the main use of EPD(M) is in outdoor applications, such as automotive sealing systems, window seals and roof sheeting, and in under-the-hood applications, such as coolant hoses. The main drawback of EPDM is its poor resistance to swelling in apolar fluids such as oil, making it inferior to high-performance elastomers, such as fluoro, acrylate and silicone elastomers in that respect. Over the last decade thermoplastic vulcanisates, produced via dynamic vulcanisation of blends of polypropylene (PP) and EPDM, have been commercialised, combining thermoplastic processability with rubber elasticity [8, 9]. EPDM was developed and commercialised in the late 1950s. With an annual production capacity of more than 1,000 kt in 1998 [10]. EPDM is currently the fourth elastomer by volume and has become more or less a commodity rubber. Actually, EPDM is the largest non-tyre rubber. The annual growth rate is about 4%. DSM and Exxon are market leaders with a combined market share of approximately 40%. PP/EPDM-based thermoplastic vulcanisates which have currently the fastest growing rubber market (8% per year).

6.1.2 EPDM Crosslinking [1-3] Most EPDM applications require crosslinking except when used as an impact modifier for PP, polystyrene (PS) and polyamides or as an oil additive, e.g., as viscosity index improver or dispersant. Most commonly, accelerated sulfur vulcanisation is used for the crosslinking of EPDM. As a result of the low amount of unsaturation in EPDM (< 1 mole/ kg versus NR - 15 mole/kg), sulfur vulcanisation of EPDM is rather slow and a relatively large amount of accelerators is needed. Because of the low polarity of EPDM the solubility of polar accelerators is limited, often resulting in low effectivity and/or blooming. Typically, up to 5 different accelerators are used in EPDM formulations. As for other rubbers environmental issues, such as nitrosamine formation and may be in the future the presence of zinc, are prompting the development of new accelerator systems.

208

Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy Saturated EPM can be cured with peroxide/co-agent combinations, but EPDM is cured even more efficiently due to the presence of the residual unsaturation. The share of peroxidecured EPDM is gradually increasing at the expense of sulfur-vulcanised EPDM [11] because of the high thermal stability of the carbon-carbon crosslinks formed upon peroxide-curing of EPDM in combination with the increasing high temperature demands in the market. Peroxide-cured EPDM has a superior compression set at temperatures above 100 °C and it is generally more heat-resistant than that of sulfur-vulcanised EPDM. In addition, peroxidecuring is not affected by environmental issues, such as nitrosamine formation and use of large amounts of the heavy metal zinc. However, with respect to ultimate and dynamic properties sulfur-vulcanised EPDM is superior to peroxide-cured EPDM as a result of the dynamic rearrangement of sulfidic crosslinks during sulfur vulcanisation, yielding a rubber network which is relatively free of internal stresses [12, 13]. Resols (phenol-formaldehyde resins) are commercially used for effective crosslinking of EPDM in the production of thermoplastic vulcanisates [8]. General studies on rubber crosslinking for different diene rubbers are presented here.

6.1.3 Studies into the Chemistry of Rubber Crosslinking The technology of sulfur vulcanisation of unsaturated elastomers has evolved since Goodyear’s invention in 1839. Scientific studies into the chemistry of sulfur vulcanisation started to appear in the late 1950s (for reviews see References [14-18]). Two experimental approaches can be distinguished: the analysis of rubber vulcanisates themselves and the so-called low-molecular-weight model studies. Analysis of rubber vulcanisates is a direct approach and is, therefore, to be preferred. Initially, it suffered from several drawbacks. The number of structures formed during sulfur vulcanisation is rather large. For example crosslinks with variable sulfur bridge length, crosslink precursors (pendent sulfur) and cyclic sulfides, and the corresponding concentrations are small relative to the detection limits of classical spectroscopic techniques, such as infra-red (IR) and Raman spectroscopy. In the case of Raman spectroscopy, fluorescence is encountered as a result of the presence of the vulcanisation chemicals [19]. The presence of fillers, especially carbon black, which are strongly lightscattering and/or light-absorbing, results in a (strongly) decreased sensitivity for optical spectroscopy methods in general. Finally, a rubber vulcanisate consists of an infinite network of crosslinked polymer chains, rendering it insoluble and, thus, inaccessible to chromatographic techniques and liquid-state nuclear magnetic resonance (NMR). Therefore most progress towards the understanding of sulfur-vulcanisation chemistry was originally made by ‘vulcanising’ low-molecular-weight model olefins. Numerous

209

Spectroscopy of Rubbers and Rubbery Materials model compound studies by the Tun Abdul Razak Laboratory (UK), the Dunlop Research Centre (Canada), the Experimental Station of E.I. DuPont de Nemours & Co. (USA), DSM Research (The Netherlands) and the Universities of Leiden (The Netherlands), South Elisabeth (South Africa), and Cologne (Germany), among others, have been published, but a comprehensive review of this work is still not fully available. Because of their low-molecular weight, the corresponding model crosslinks are soluble and sometimes even volatile, allowing chromatographic techniques, such as gas chromatography and liquid chromatography, for separation and also liquid-state NMR for identification. Clearly, the results of model olefin studies should be used with care. Several practical pitfalls have been encountered in the past and a ‘translation’ of the results to the polymer system may also be a source of errors [20, 21]. Over the last decade the development of advanced analytical techniques, such as Fourier transform (FT) Raman and solid-state NMR spectroscopy, have been impressive, resulting in a great deal of progress in the field of the sulfur vulcanisation of unsaturated elastomers [22-25]. In general, most of the problems encountered in the study of the chemistry of the sulfur vulcanisation of elastomers are also encountered in the study of peroxide-curing. In comparison with sulfur vulcanisation only a limited number of spectroscopic studies on peroxide-curing have been published.

6.1.4 Scope The aim of this chapter is to review optical spectroscopy studies on sulfur and peroxide crosslinking of polydiene rubbers, such as NR and BR (Sections 6.2.1 and 6.3.1, respectively), and to discuss in detail recent FT-Raman and FT-IR spectroscopy studies into the sulfur and peroxide crosslinking of EPDM (Sections 6.2.2 and 6.3.2, respectively). The results of optical spectroscopy studies will also be discussed in the light of results obtained with other techniques. Finally, the elucidation of the chemical structures of the crosslinks formed will allow enhanced understanding of the mechanisms of crosslinking and some preliminary insight into the structure/property relationships of crosslinked rubber. Raman and IR spectroscopic studies dealing with the qualitative and/or quantitative determination of rubber compounding ingredients, i.e., the elastomer itself [22, 26-31], fillers [32, 33], vulcanisation chemicals and other additives [34-37], are not included here. The same applies to studies dealing with the crosslinking of elastomers by means of chemicals other than sulfur or peroxide [38-41], self-crosslinking of elastomers blends [42-44], crystallisation (strain-induced) [45-48] and oxidation/ageing [49-53].

210

Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy Raman and IR spectroscopy are similar techniques, since both are used to study oscillations of structural moieties of molecules, namely stretching and bending vibrations [54]. Raman and IR spectroscopy are different in the sense that strong Raman bands are associated with vibrational motions of highly polarisable bonds and, so, Raman spectroscopy is more sensitive for vibrations of apolar groups. In contrast, intense IR bands are primarily due to motions of bonds with strong permanent dipole moments and thus, IR spectroscopy is more sensitive for vibrations of polar groups. As a result, Raman and IR spectroscopy are usually seen as complementary techniques which have proven their value in polymer science and technology [55-56]. FT-Raman spectroscopy was used for studying the sulfur vulcanisation of EPDM, since apolar moieties, such as the residual unsaturation of the third monomer of EPDM, and sulfur crosslinks are involved (Section 6.2.2). FT-IR spectroscopy was used to study the contribution of polar co-agents during peroxide-curing of EPDM (Section 6.3.2). The FT technique was essential for both studies to obtain high-quality spectra with a sufficient signal-to-noise ratio. For Raman spectroscopy the problem of fluorescence was overcome by excitation of the Raman scattering at 1064 nm using a Nd:YAG laser.

6.2 Sulfur Vulcanisation

6.2.1 Sulfur Vulcanisation of Polydiene Rubbers

6.2.1.1 General In the late 1960s optical spectroscopy was first mentioned for studies on the sulfur vulcanisation of unfilled rubber. For instance, Stewart and Linnig showed, using IR spectroscopy, that sulfur vulcanisation of NR in the presence of tetramethyl thiuram disulfide (TMTD) and ZnO resulted in the formation of zinc dimethyl dithiocarbamate (ZDMC) [57]. Some years later three series of optical spectroscopy studies were published by the Koenig group [19, 58-63], which can be seen as a breakthrough, since they provided the first structural data for sulfur vulcanisation of high-molecular-weight elastomers. These studies allowed the verification of (intermediate) structures identified in the series of low-molecular-weight studies by Porter and co-workers and the validation of the mechanisms proposed [14, 16]. However, several experimental problems had to be overcome before final conclusions could be drawn from Raman spectra. Firstly, the problem of strong fluorescence due to the presence of accelerator impurities and/or

211

Spectroscopy of Rubbers and Rubbery Materials oxidation products, which resulted in low signal/noise spectra, had to be solved by choosing optimum vulcanisation conditions, carefully extracting the vulcanisates and laser irradiation of the samples for several hours before recording the spectra [19, 58]. Secondly, because of a strong overlap of the various peaks, the Raman spectra had to be curve-resolved [58, 59]. Finally, tentative assignments had to be verified by using selective extraction solvents and/or chemical probes for the selective conversion of disulfides and polysulfides [59].

6.2.1.2 Sulfur Vulcanisation of cis-BR Table 6.1 gives an overview of the structures that were finally identified by Koenig and co-workers using Raman spectroscopy during the vulcanisation of cis-BR at 140 or 150 °C [19, 58-59]. For sulfur-free vulcanisation of BR using only TMTD and ZnO not only the formation of ZDMC but also the formation of alkenylsulfides and of monosulfidic, disulfidic and polysulfidic crosslinks were demonstrated. In addition, main-chain modifications were observed, such as cis-trans isomerisation and the formation of conjugated trienes and cyclic unsaturated sulfides. For sulfur vulcanisation of cis-BR in the presence of mercaptobenzothiazole (MBT), ZnO and lauric acid, similar structures were identified with the exception of monosulfides. Sulfur/TMTD/ZnO/lauric acid vulcanisation was studied as a function of time and the results were interpreted in combination with data obtained from crosslink density measurements. As expected, similar kinetics were found for the formation of alkenylsulfides and sulfidic crosslinks as well as the development of the crosslink density, until a plateau was reached. Interestingly, the triene formation followed a similar time profile. However, cis-trans isomerisation and diene formation were shown to continue beyond the optimum cure time, indicating that they were independent of the vulcanisation process. Using IR the degree of cis-trans isomerisation of cis-BR during sulfur vulcanisation was quantified. Bishop showed high degrees of isomerisation from 95% to about 35% cis if large amounts of sulfur were used for vulcanisation [64], whereas Madge [65] and Blümel [66] found only modest degrees of isomerisation from 95% to 88%, if much lower sulfur concentrations were applied. These high levels of isomerisation again demonstrate that isomerisation and vulcanisation are independent processes. Devlin showed that in the presence of carbon black the cis-trans isomerisation was markedly reduced from 97% to 94% using ‘typical’ amounts of sulfur [67]. Coleman of the Koenig group identified the reaction products obtained after heating various mixtures of sulfur, TMTD, tetramethyl thiurammonosulfide, ZDMC and/or ZnO in the absence of elastomer or olefin using Raman spectroscopy. This enabled the

212

Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy

Table 6.1 Raman identification of structures formed upon accelerated sulfur vulcanisation of cis-BR [19, 58, 59] Raman band (cm-1)

Structure

1575 to 1700

C=C stretching

1664

trans C=C

1650

cis C=C

1640 1633

vinyl C=C C=C-C-S

1605

conjugated C=C-C=C

1623

conjugated C=C-C=C-C=C

1587

conjugated C=C-C=C-C=C

400 to 900

C-S and S-S stretching

744

alkenyl monosulfide C-S

720

alkenyl disulfide C-S

696

cyclic sulfide C-S

635

unsaturated 6-ring cyclic sulfide C-S

577 505 - 510

pendent dithiocarbamate disulfide S-S

505

unsaturated 5-ring cyclic sulfide C-S

424

polysulfide S-S

a

a: tentative assignment

elucidation of the complex thiuram sulfide and dithiocarbamate chemistry [60, 61]. These reactions will probably also occur during sulfur vulcanisation of unsaturated elastomer, but it has to be realised that the reaction mechanism in the latter case is more intricate and selectivities may be totally different. Intermediate free radicals, such as the thiuram radicals and the thiurampersulfenyl radicals, were identified using electron spin resonance spectroscopy. Sulfur and/or TMTD vulcanisation of BR proceed probably via free radical mechanisms, whereas in the presence of ZnO or ZDMC contributions of ionic pathways will dominate.

213


6.2.1.3 Sulfur Vulcanisation of NR Chen and co-workers tentatively assigned new signals in the FT-IR spectra of accelerated sulfur-vulcanised NR to the formation of C-S and S-S bonds corresponding to monosulfides, disulfides and polysulfides [68]. The vulcanisation of NR was retarded when clay was added to the NR compound. Recently, it was shown using FT-IR that the decrease of vulcanisation rate and final crosslink density of sulfur vulcanised NR upon increasing silica content may be related to increased absorption of zinc stearate onto the silica surface [69]. In a third series of papers from the Koenig group, reversion of accelerated sulfur-vulcanised NR was studied using IR spectroscopy in combination with crosslink density measurements [62, 63]. Chen and co-workers showed that at 140 to 150 °C reversion of sulfur-vulcanised unfilled NR occurs at high sulfur/accelerator ratios, which could be related to the isomerisation of the cis-2-methyl-2-butene monomeric unit of NR into a trans-2-methyl-3-butene unit and to the formation of conjugated trienes. It was shown that reversion is a result of neither oxidation nor isomerisation of the cis-2-methyl-2butene unit into the corresponding trans unit. When NR compounds were extended with carbon black, FT-IR spectra could still be recorded for carbon black levels as high as 30 phr. In the presence of carbon black less cis-trans isomerisation was observed, which was explained by the higher thermal stability of polysulfides as a result of polymerfiller interactions. It was also shown that the reversion resistance was improved if more reinforcing carbon black was used. Schotman and co-workers tentatively assigned the new Raman peaks at 1625 and 1592 cm-1 observed during sulfur vulcanisation of squalene, to the formation of conjugated dienes and trienes, respectively [70]. When vulcanisation was carried out in the presence of 1,3-di(citraconimidomethyl)benzene, this resulted in a reduced intensity of these two new peaks, corroborating that conjugated dienes and trienes, formed as a result of reversion, react with the diimide. Obviously, the diimide is not an anti-reversion agent in the sense that it prevents reversion, but it is in the sense that it repairs crosslinks when reversion has occurred.

6.2.1.4 Mechanism of Sulfur Vulcanisation of Polydiene Rubbers The results of the optical spectroscopy studies into sulfur vulcanisation of polydiene rubbers correspond well with the results obtained via low molecular weight model olefin studies and solid state 13C NMR studies. From all these studies the mechanism for accelerated sulfur vulcanisation as shown in Figure 6.2 has emerged [14-18], which is

214


Figure 6.2 Simplified scheme of the mechanism of accelerated sulphur vulcanization of polydiene rubbers [14-18]

now commonly accepted, although it should be realised that some details of the various reaction steps still need to be refined. First, sulfur, accelerator and activators react towards a soluble sulfurated zinc complex (activated sulfur). This complex reacts with the unsaturated elastomer by substituting a labile allylic hydrogen atom, which results in the attachment of accelerator residues to the elastomer chain (pendent sulfur or crosslink precursor). This intermediate is converted into a sulfur crosslink either via disproportionation with a second pendent sulfur structure or by allylic substitution of the unsaturation of a second elastomer chain. Recently, it was shown that the former route is probably most likely [71]. It is noted that during accelerated sulfur vulcanisation dialkenylsulfides are predominantly formed, indicating that the unsaturation is not consumed during reaction but that it activates the α-position, whereas during unaccelerated sulfur vulcanisation mixtures of dialkenylsulfides, dialkylsulfides and alkenylalkylsulfides are formed. Upon prolonged vulcanisation

215

Spectroscopy of Rubbers and Rubbery Materials network maturation reactions may occur, for example, the sulfidic crosslink may exude sulfur, resulting in crosslink shortening, and/or it may be converted into cyclic sulfides, which is usually associated with reversion.

6.2.2 Sulfur Vulcanisation of EPDM

6.2.2.1 Early Studies The mechanism of the accelerated sulfur vulcanisation of EPDM is probably similar to that of the highly unsaturated polydiene rubbers. The vulcanisation of EPDM has been studied with emphasis on the cure behaviour and mechanical and elastic properties of the crosslinked EPDM. Hardly any spectroscopic studies on the crosslinking chemistry of EPDM have been published, not only because of the problems discussed in Section 6.1.3 but also because of the low amount of unsaturation of EPDM relative to the sensitivity of the analytical techniques. For instance, high-temperature magic-angle spinning solidstate 13C NMR spectroscopy of crosslinked EPDM just allows the identification of the rubber type, but spectroscopic evidence for the presence of crosslinks is not found [72]. Only two spectroscopic studies on sulfur vulcanisation of EPDM by Fujimoto and coworkers are available [73-74]. Using attenuated total reflectance (ATR) IR spectroscopy they showed that during sulfur/TMTD/MBT/ZnO/stearic acid vulcanization, the C=C bands at 3035, 966 and 870 cm-1 of the residual unsaturations of the EPDM third monomers, DCPD, 1,4-hexadiene (HD) and 5-methylidene-2-norbornene (MNB), respectively, decreased in intensity as a function of time at 140 and 150 °C. The relative decrease in intensity was shown to correlate with the increase in crosslink density. In Sections 6.2.2.2 and 6.2.2.3 it will be shown that this decrease of intensity should not be interpreted as a loss of unsaturation during sulfur vulcanisation of EPDM. A series of low-molecular-weight model studies using 2-ethylidene norbornane (ENBH) by Duynstee and co-workers have provided a lot of detailed structural information on the sulfur vulcanisation of EPDM [75, 76]. As for the polydiene rubbers, accelerated sulfur vulcanisation of EPDM proceeds via allylic substitution, resulting in the formation of dialkenylsulfides. The number of ENBH model crosslinks that has been separated by preparative high-pressure liquid chromatography (HPLC) and identified with 1H NMR amounts to about 40 as a result of the variation of the sulfur bridge length (1 to 5 sulfur atoms), the Entgegen/Zusammen isomerism of the ENB ethylidene unit and the 3 different allylic positions for sulfur attachment for both ENBH moieties (C3,exo, C3,endo and C9). One of these ENBH-Sn-ENBH model crosslink structures is shown in Figure 6.3.

216


Figure 6.3 Accelerated sulphur vulcanization of EPDM as demonstrated by lowmolecular weight model olefin studies [75,76] After these early studies an extensive FT-Raman study [77] was performed to bridge the gap between the low-molecular-weight ENBH model vulcanisation studies and the vulcanisation studies using high-molecular-weight EPDM. These studies will be presented in detail. First, a series of low-molecular-weight dialkenylsulfides will be discussed in order to determine the effect of sulfur vulcanisation on Raman spectra of olefins. Subsequently, the attachment of the sulfur crosslinks at the allylic positions, the conversion of ENB, the length of sulfur crosslinks and the network structure will be addressed for unfilled sulfur vulcanisates of amorphous EPDM. Some preliminary network structure/ properties relationships will also be presented.

6.2.2.2 Dialkenylsulfides It has been shown that during the accelerated sulfur vulcanisation of unsaturated elastomers or model olefins allylic hydrogen atoms are substituted by sulfur atoms (Figures 6.2 and 6.3). So, the C=C vibrations at 1600 - 1800 cm-1, the S-S peaks at 300 - 700 cm-1 and the C-S bands at 600 - 750 cm-1 provide the most relevant information for monitoring sulfur vulcanisation using Raman spectroscopy. Table 6.2 gives FT-Raman data for a series of alkenes and the corresponding dialkenylsulfides. The latter were either synthesised as pure sulfides [78] or produced as a mixture via model olefin vulcanisation [79]. Sulfur substitution of the allylic hydrogen atoms results in a downward shift of the wave number of the C=C stretching vibration of 7 to 18 cm-1. For 3-methyl-1-pentene an upward shift of 20 cm-1 is observed, but this is due to complete allylic rearrangement of the unsaturation towards a more highly substituted unsaturation upon vulcanisation [80]. When compared with 3methyl-2-pentene a downward shift of 13 cm-1 is observed. The S-S vibrations have been assigned using the spectra of pure di(2,3-dimethyl-2butenyl)disulfide, trisulfide and tetrasulfide. Literature data on dialkenylsulfides are available only for diallyldisulfide and disqualenylsulfides (Table 6.2), but more data for

217


Table 6.2 T-Raman spectral data of alkenes and (mixtures of) corresponding dialkenylsulfides FT-Raman band (cm-1)

Dialkenylsulfide alkene

sulfide

alkene

dialkenylsulfide

C=C

C=C

S-Sa

mono

1675

1666

-

di

1675

1663

510

tri

1675

1663

490

tetra

1675

1662

491/439

mixture

1675

1661

490/438/459

2-methyl-2-pentene

mixture

1676

1663

475/460/494/439

3-methyl-1-pentene

mixture

1640 (1673b)

1660

491/460/439

3-hexene

mixture

1671

1660

459/437/503/475

ENBH

mixture

1688

1678

492/441

DCDPH

mixture

1612

1605

484/450/436

allene [19, 81]

di

-

-

510

squalene [125]

mixture

1667

1649

2,3-dimethyl-2-butene

a: when two or more wave numbers are given, the first one corresponds to the Raman peak with the largest intensity; b: due to complete allylic rearrangement of 3-methyl-1-pentene upon sulfur vulcanisation, the comparison should be made with 3-methyl-2-pentene

dialkylsulfides are available by Freeman [81]. For dialkyldisulfides and dialkyltrisulfides, one peak each is observed in the ranges of 500 to 520 cm-1 and 480 to 500 cm-1, respectively. For dialkyltetrasulfides two peaks are found at about 440 and 490 cm-1. Dialkylpentasulfides are characterised by a peak at 490 cm-1 and two smaller peaks at 460 and 440 cm-1. The observation of more than one peak for the higher sulfides is due to the presence of different S-S conformations. A comparison of the data for dialkenylsulfides (Table 6.2) and the literature data for dialkylsulfides [81] shows that the nature of the substituents on the sulfide (alkenyl versus alkyl) does not affect the position of the Raman bands.

218

Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy The FT-Raman spectra of the sulfur vulcanisates of the various model olefins do not contain the characteristic disulfide signal at 510 cm-1, but do contain the typical higher sulfide bands at 490, 460 and 440 cm-1 (Table 6.2). In addition, a new band at about 475 cm-1 is observed for the vulcanisates of 2-methyl-2-pentene and 3-hexene, which has not yet been assigned (hexasulfide?). Results of HPLC analysis have shown that the vulcanisate of 2,3-dimethyl-2-butene consists mainly of a mixture of disulfide to pentasulfide with about 15 mole% of disulfide [79]. This illustrates that FT-Raman spectroscopy is not very sensitive for the identification of disulfides. Because of an overlap of signals, FT-Raman does not provide detailed, quantitative information on the presence of the individual higher sulfides (S>2).

6.2.2.3 Unfilled Vulcanised EPDM

6.2.2.3.1 ENB Conversion It has been shown that during sulfur vulcanisation of EPDM the C=C peak of the residual ENB unsaturation at 1685 cm-1 seems to decrease in intensity in agreement with the observations by Fujimoto and co-workers [73, 74] (see Section 6.2.2.1). However, in Section 6.2.2.2 it was shown that sulfur vulcanisation of the low-molecular-weight ENBH results in a shift of the Raman C=C peak from 1688 to 1678 cm-1. Taking this into account a closer inspection of the FT-Raman spectra reveals that the original C=C peak at 1690 cm-1 decreases in intensity, and a new peak is observed at 1681 cm-1. Actually, the C=C peak broadens towards lower wave numbers, but in a first approximation the total area remains constant. So, the sulfur substitution reaction of the allylic hydrogens is confirmed for the polymer system. This corresponds to the observation by Koenig and co-workers, namely that upon sulfur vulcanisation of cis-BR, the C=C peak at 1650 cm -1 decreases in intensity and that of a new peak at 1633 cm-1 increases its intensity [19, 58]. The intensity changes in the FT-Raman spectra of EPDM upon vulcanisation have been used to quantify the chemical conversion of this reaction. The C=C stretching region was fitted for a series of spectra using a 3-dimensional fitting program, fitting identical band positions and widths for all spectra, fitting the intensities of the bands individually for each spectrum [77]. The two bands were set at 1690 cm-1 (unreacted ENB) and 1681 cm-1 (reacted ENB). The line width of the C=C vibration increases from 11 to 21 cm-1 upon sulfur vulcanisation. This increase is larger than that for the low-molecular-weight model ENBH (10 to 13 cm-1), which suggests that more lower sulfides may be formed for EPDM. By using the 1450 cm-1 cluster as a reference it was determined that the Raman scattering cross-section of the 1681 cm-1 band is 30% higher than that of the 1690 cm-1 band and so, the conversion of ENB can be calculated (Table 6.3; the NMR data in this

219

220

46

63

63

65

65

65

65

65

8.9

8.6

4.6

10.0

10.0

10.0

10.0

10.0 3.0

1.5

1.2

85

40

34

25

14

0.6b 0.9

63

43

48

41

1.5

1.5

1.5

1.5a

0.71

0.33

0.28

0.21

0.12

0.24

0.31

0.36

0.30

1. 6

2.7

3.4

4.6

7.8

3.1

3. 1

3.0

3.4

molecular weight between crosslinks (kg/mole)

NMR

0.31

0.18

0.15

0.11

0.064

0.16

0.16

0.17

0.15

crosslink density (mole/kg)

2.3

1.8

1. 9

1, 9

1.8

1.5

1. 9

2.1

2.0

crosslink stoichiometry (mole/mole)

FT-Raman + NMR

a: EPDM compounds containing 5 phr ZnO, 1 phr stearic acid, 1 phr TMTD, 0.5 phr MBT and 1.5 phr sulfur were vulcanised for 30 minutes at 160 °C; b: for the compounds with different sulfur concentration the concentrations of ZnO, stearic acid, TMTD and MBT were adjusted proportionally

33

8.9

ENB converted ENB conversion (mole/kg (%) EPDM)

sulfur content (phr)

ENB content (wt.%)

Mooney viscosity ML(1+4)125°C

FT-Raman

Sample

Table 6.3 Number of converted ENB units as determined by FT-Raman spectroscopy and number of chemical crosslinks as determined by solid state 1H NMR relaxation studies of unfilled EPDM vulcanisates


Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy table will be discussed in Section 6.2.2.3.3). The number of ENB molecules reacted with sulfur can be calculated from the fraction of ENB converted, as determined by FT-Raman spectroscopy, and the composition of the vulcanisation recipe. From the data in Table 6.3 it follows that the relative ENB conversion ranges from 15% to 85%, which is larger than often assumed. For the EPDM vulcanisates with 1.5 phr of sulfur the absolute ENB conversion is about 300 to 350 mmole ENB/kg EPDM. This conversion is not affected much by the Mooney viscosity of the original EPDM nor by the original ENB content of the EPDM. The latter indicates that for these samples it is not the ENB content but some other parameter is limiting the ENB conversion. When the Raman data for the vulcanisates of the EPDM with 10.0 wt.% ENB and sulfur contents varying from 0.6 to 3 phr are plotted against the sulfur content a straight line going through the origin is obtained (Figure 6.4). Toluene extraction of these vulcanisates followed by thin layer chromatography quantification showed that the amount of free sulfur (S8) is less than 1% of the original content. Since the amount of sulfur present in loop/cyclic structures is small (< 10%) (compare with Section 6.2.2.3.3), it is concluded that the sulfur added is predominantly used for crosslinking. It can be calculated that the average length of the sulfur crosslink is 2.7 sulfur atoms for the ISO 4097 [82] formulations studied. For a given ratio of sulfur to accelerators the average length of the sulfur crosslink seems to be constant, independent of the EPDM type used.

Figure 6.4 Absolute ENB conversion as determined with FT-Raman spectroscopy versus sulphur content of EPDM compounds (data from Table 6.3)

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6.2.2.3.2 Sulfur Crosslinks The region between 300 and 700 cm-1 in FT-Raman spectra contains information regarding the nature of the sulfur crosslinks. From the studies of low-molecular-weight sulfides it is concluded that the location of bands assigned to disulfides, trisulfides and tetrasulfides is at approximately 500-520, 485-500 and 490 + 440 cm-1, respectively [81]. In the FTRaman spectrum of sulfur vulcanised EPDM it can clearly be seen that the bands are present at 570, 520, 490 and 440 cm-1. Acetone extraction experiments showed that the peak at 570 cm-1 is due to the presence of ZDMC, formed from ZnO and TMTD. ZDMC also has a major band at 440 cm-1, which coincides with the tetrasulfide band at the same wave number. The 490 cm-1 band results from trisulfides and higher sulfides. The 520 cm-1 band is assigned to disulfides, which was not observed for the low-molecularweight model olefin vulcanisates (compare with Section 6.2.2.2). Although the average length per sulfur bridge is only 2.7 sulfur atoms, the relative intensity of the tetrasulfide bands is still considerable. However, one has to realise that the polarisability increases as the sulfur rank of the crosslink increases. Consequently, the Raman scattering crosssection rapidly increases with the length of the sulfur bridge. Although detailed analysis of the 400-700 cm-1 region in the FT-Raman spectrum of sulfur vulcanised EPDM is still a problem due to the poor signal:noise ratio, it appears that the relative intensities of the bands in this region do not change with the conversion of the third monomer or the type of EPDM. The trisulfide band at 490 cm-1 can be observed without the problem of overlap. In the series of EPDM with 10 wt.% ENB, with increasing amounts of vulcanisation chemicals the area of this band increases, which is in agreement with the increase in ENB conversion observed. The fact that the relative amounts of disulfide, trisulfide and tetrasulfide peaks are not affected can be explained by the fact that for all samples, the relative amounts of the various vulcanisation ingredients is kept constant. Using chemical probe treatments [83] attempts were made to support the assignment of the Raman sulfide bands for sulfur-vulcanised EPDM. Treatment of the vulcanisate based on EPDM with 10 wt.% ENB and 1.5 phr sulfur with 1-hexanethiol/piperidine results in the disappearance of all signals in the 300 to 600 cm-1 region. Treatment with 2propanethiol/piperidine/hexane results in the disappearance of the 490 cm-1 band, but the 440 cm-1 peak remains and the 510 cm-1 signal actually increases in intensity. The various reactions of sulfur crosslinks possibly occurring in the presence of alkylthiols are depicted in Figure 6.5. Alkanethiols are not able to convert monosulfides (I), indicating that C-S bonds are not affected [83]. Therefore, reactions II and VI will not occur either. For low-molecular dialkenylsulfides it has been demonstrated that 1-hexanethiol degrades all disulfidic and higher sulfidic species. According to reaction III, a disulfide crosslink is converted into R-SH and R-S2-X. Since the applied molar ratio, 1-hexanethiol/sulfur

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Figure 6.5 Possible reactions of sulphur crosslinks R-Sm-R (R = rubber) with thiols XSH (X= alkyl) crosslink is larger than 100, R-S2-X will be converted according to reaction IV, which is similar to reaction III. The combination of reactions III and IV will result in the ultimate formation of R-SH and X-S2-X. R-SH rubber thiols do not contain S-S bonds and X-S2X is washed from the rubber, all sulfidic bands in the 300-600 cm-1 region have disappeared. Experiments with low-molecular-weight dialkenylsulfides have demonstrated that treatment with 2-propylthiol results in the selective degradation of trisulfides and higher sulfides probably for electronic reasons [83]. Reactions VI and VII are the first reactions to occur. In a reaction sequence as discussed for 1-hexanethiol all trisulfides and higher sulfides will finally be degraded to R-S2-X, R-S2H, R-SH and X-S2-X. The Raman spectrum of the 2-propanethiol-treated EPDM vulcanisate indeed shows that the trisulfide and higher sulfides peak at 490 cm-1 has disappeared. In addition, the intensity of the disulfide peak at 510 cm-1 has increased. Since hexane is used as a solvent, ZDMC will probably not dissolve and therefore, the observation of the band at 440 cm-1 is related to the presence of this chemical. Tetrasulfides and higher sulfides do not contribute to this signal, since the characteristic peak at 490 cm-1 is absent. So, it is concluded that the chemical probe treatment has confirmed the assignment of the various sulfidic crosslinks formed during EPDM vulcanisation.

6.2.2.3.3 Network Structure NMR relaxation time measurements provide information on the mobility of polymer chain segments and have been used for studying the effect of crosslinking, oil extension

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Spectroscopy of Rubbers and Rubbery Materials and reinforcement of rubbers. Recently, a wideline solid-state 1H NMR method was developed for determining the spin-spin 1H relaxation time, which can be converted into the molecular weight between chemical crosslinks in rubbery networks (Mc) and, thus, into the density of chemical crosslinks (1/2Mc) [84]. Using this approach the crosslink density of the sulfur-vulcanised EPDM samples studied with FT-Raman spectroscopy could be confirmed (Table 6.3). In agreement with the FT-Raman data, but with small experimental error, the NMR data show that for all unfilled EPDM vulcanisates with 1.5 phr sulfur the number of chemical crosslinks is more or less constant independent of the elastomer type, i.e., 160-180 mmole/kg. For all samples, except the vulcanisate of EPDM with 4.6 wt.% ENB and 1.5 phr sulfur, the ratio of the converted number of ENB moieties and the number of chemical crosslinks is around 2.0 (± 0.2). The statistical spread in the ratio of the converted number of ENB molecules and the number of chemical crosslinks is estimated to be 10-15% based on the error in the Raman and NMR data. Furthermore, there might be a systematic error in the crosslink density as determined by NMR, originating from the assumptions made in calculating the number of chemical crosslinks from the total number of crosslinks (chemical and entanglements). The experimental ratios of 2 (± 0.2) agree with the expected value of 2.0 converted ENB units for one chemical crosslink (ENB-Sn-ENB), indicating that crosslinks are formed predominantly. Several explanations are proposed for deviations from the theoretical value. Both pendent sulfur as in the crosslink precursor (ENB-Sn-X, X = accelerator residue) and sulfur loops (sulfur bridge between two ENB units in the same EPDM chain) will give rise to a ratio larger than 2.0. Formation of cyclic sulfide (addition of S to unsaturation) will give rise to a ratio lower than 2.0. The fact that within experimental error a ratio of 2.0 is found suggests that pendent sulfur, loops and cyclic sulfur hardly occur (less than 10% of the total amount of reacted sulfur).

6.2.2.3.4 Network Structure/Properties Relationships The analytical techniques discussed previously can be used to study the EPDM network as such or its formation in time as well as to determine relationships between the network structure and the properties of the vulcanisates. In a preliminary approach some typical vulcanised EPDM properties, i.e., hardness, tensile strength, elongation at break and tear strength, have been plotted as a function of chemical crosslink density (Figure 6.6). The latter is either determined directly via 1H NMR relaxation time measurements or calculated from the FT-Raman ENB conversion (Table 6.3). It is concluded that for these unfilled, sulfur-vulcanised, amorphous EPDM, the chemical crosslink density is the main parameter determining the vulcanisate properties. It is beyond the purpose of this review to discuss these relationships in a more detailed and theoretical way.

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Figure 6.6 Mechanical properties of unfilled sulphur vulcanized amorphous EPDM as a function of crosslink density

6.3 Peroxide-curing

6.3.1 General The mechanism of peroxide crosslinking of elastomers is much less intricate than that of sulfur vulcanisation. Crosslinking is initiated by the thermal decomposition of a peroxide, which is the overall cure rate determining step. Next, the active radicals thus formed abstract hydrogen from elastomer chains to form macroradicals. Finally, crosslinking results either from the combination of two macroradicals or from the addition of a macroradical to an unsaturated moiety of another primary elastomer chain. In view of the abundance of unsaturation in polydiene elastomers it may be expected that this type of elastomers can be cured very efficiently with peroxides. This section will show that a very high crosslinking efficiency (≡ moles of crosslinks formed per mole of peroxide decomposed) can indeed be obtained. However, the substitution pattern of the unsaturation plays a major role in the actual efficiency found. The industrial relevance of peroxide-curing of elastomers is by far larger for main-chain saturated elastomers, such as silicone elastomers, acrylic elastomers, fluoro elastomers

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Spectroscopy of Rubbers and Rubbery Materials and particularly EPDM. Nonetheless, there has been little optical spectroscopical work studying the curing kinetics and mechanism. The EPDM studies that have been published indicate that detailed information on the curing chemistry can be obtained. The role of the type and the amount of the third monomer in particular, as well as the role of coagents and scorch retarders, have been elucidated.

6.3.2 Polydiene Elastomers 6.3.2.1 Natural Rubber In early communications it was reported that the efficiency of dicumylperoxide (DCP) crosslinking of NR amounts to unity [85, 86], which means that one mole of chemical crosslink is obtained as the result of the thermal decomposition of one mole of DCP. This was explained by assuming that crosslinking results from the combination of polyisopropenyl radicals only and crosslinking by macroradical addition does not take place. It is not easy to support this crosslinking mechanism with spectroscopic evidence, considering that the differences in peak intensity of the =C-H out-of-plane bending mode, absorbing at 838 cm-1, before and after peroxide-curing, are difficult to quantify. Moreover, the formation of crosslink structures is even more difficult to detect, because C-C stretching modes have a very weak intensity. Nonetheless, data processing techniques, such as the least squares curve-fitting procedure as developed for polymeric systems by Koenig and co-workers [87], have proven their value in this respect. Via this method, the spectrum of a mixture (in this case NR, DCP and the peroxide decomposition products) could be fitted with the spectra of the pure components and the fractional amount of each component was calculated, assuming that the absorptions of the various components were additive in the mixture. For NR cured with an excessive amount of DCP (> 50 phr!) it was shown that a decrease of double-bond concentration occurred at a rate comparable with the formation of peroxide decomposition products. This indicates that the crosslinking efficiency can exceed unity as opposed to the earlier reports [85, 86] mentioned. FT-IR results did not provide evidence for cis-trans isomerisation, although isomerisation was confirmed from solid-state NMR data that were generated on the same samples. Interestingly, the formation of quaternary carbon-carbon bonds resulting from crosslinking could be detected at 1320 cm-1, but quantification was impossible in the light of the weak intensity and width of the peak. In more recent studies from Gonzalez and co-workers [88-90] it was concluded from dynamic mechanical analysis of peroxide-cured NR that a non-uniform crosslinked network results if a large amount of peroxide is used. This result seems to be in line with the optical spectroscopy studies discussed.

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6.3.2.2 Polybutadienes Patterson monitored the peroxide-curing behaviour of cis-BR by both FT-IR and FTNMR techniques, again using excessive amounts of DCP [87]. It was concluded that cistrans isomerisation occurred, based on a new absorption appearing in the IR spectrum at 965 cm-1 (=C-H, out-of-plane). Moreover, it was concluded that crosslinking occurred through radical addition reactions, because the peak intensity at 728 cm-1 (cis =C-H, out of plane) decreased. Also in this case the formation of C-C bonds resulting from crosslinking reactions was shown by a weak and broad absorption appearing at approximately 1320 cm-1. Based on FT-IR data of both the double-bond consumption and the formation of peroxide-decomposition products a crosslinking efficiency of approximately 1.7 was calculated. Loan used more practical amounts of peroxide and showed that the crosslinking efficiency for cis-BR with 3% of 1,2-vinyl units is about 10, based on FT-IR data on the consumption of the vinyl double bond at 911 cm-1 (=C-H, out-of-plane) [91]. Hummel and Kaiser [92] and Van der Hoff [93] found similar values. By studying BR grades with different 1,2vinyl contents Van der Hoff showed that the crosslinking efficiency becomes as high as 50 for a product with 79% 1,2-vinyl units. FT-IR data also demonstrate that during curing significant cis-trans isomerisation occurs. Bellander and co-workers [94] showed that the crosslinking efficiency of BR also depends on the pressure applied during curing. Using the change of the vinyl absorption at 911 cm-1 and the peak at 1435 cm-1 (stretching of main-chain CH2) as an internal reference, it was concluded that the crosslinking efficiency for a BR grade with 11 wt% of 1,2-vinyl units amounted to 8, if the unfilled compounded film was cured at 170 °C in a mould at a pressure of 26 MPa. The crosslinking efficiency further increased to about 12, if the pressure during curing was increased to 293 MPa. This effect was explained by assuming that the rate of radical combination reactions is lowered as a result of the reduced mobility of polymer chains under higher pressures.

6.3.3 EPDM

6.3.3.1 Effect of Third Monomer Type and Content Fujimoto and co-workers reported in 1969 on the use of a new high-temperature ATR apparatus that they constructed for studying, among other things, the role of the type of third monomer on the peroxide curing-chemistry of EPDM [73, 74]. EPDM grades containing either ENB, DCPD, HD or MNB as a third monomer were compounded with

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Spectroscopy of Rubbers and Rubbery Materials DCP and uncured films were deposited on a multi-reflectance prism. Next, the ATR cell was heated to 140 °C and spectra were recorded as a function of curing time, while the angle of incidence was fixed at 45o, which minimised loss of radiation. The decrease of the absorption relating to the third monomer pendent unsaturation was monitored at 1685, 3045, 966 and 870 cm-1 for EPDM polymers containing ENB, DCPD, HD and MNB, respectively. The absorption was normalised with the methyl absorption at 1380 cm-1 which was not affected during curing. In the case of MNB-EPDM the absorption of the pendent unsaturation rapidly decreased within 25 minutes to about 20% of its initial value and then very slowly decreased to about 15% of its initial value. Peroxide decomposition data [95] indicate that a significant amount of peroxide was still present after 25 minutes, which means that apparently two kinetic regimes exist. Although less obvious, the data presented for the consumption of the pendent unsaturation in the case of the other EPDM grades suggest the same. Although peroxide-decomposition data have to be used with great caution, it can be concluded that double bonds are consumed by radical addition reactions. One can debate whether the unsaturations are consumed by multiple radical addition reactions or via consecutive radical addition/radical transfer sequences. The latter seems most likely, considering the low tendency of alkyl radicals for addition to alkyl-substituted double bonds under these relatively mild conditions. In radical addition reactions of this kind, the stabilisation of radicals due to polar effects is negligible. Experimental studies show that the reactivity is mainly controlled by steric effects [96]. The order of reactivity MNB > DCPD ≈ ENB > HD towards radical addition reactions as found by Fujimoto and coworkers [73, 74] is in line with these considerations. Baldwin and co-workers have monitored the reactions of the unsaturation of MNB-EPDM during peroxide-curing as well [97]. To this end thin films of unfilled compound were cured with 0.6 phr of DCP at different time intervals at 170 °C and examined for IR spectroscopic changes. It was shown that 38% of the pendent unsaturation (of 2.5 wt% MNB incorporated) was consumed after complete peroxide decomposition. This means that the number of radical addition reactions per radical species initiated amounts to about 8. Baldwin and co-workers also studied the peroxide curing behaviour of other EPDMs with a variety of third monomers [97]. As expected, they found that terminal olefins as present in 5-vinyl-2-norbornene (VNB) and MNB containing EPDM provide a much higher peroxide curing efficiency as compared to internal olefins as present in for instance ENB, DCPD and HD containing EPDM. This again is in line with the earlier considerations mentioned [96]. Based on the literature reports discussed, a rough scheme for the peroxide-curing of EPDM is proposed (Figure 6.7). This scheme is identical to the mechanism that was proposed based on the results of rheometer studies [11].

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Figure 6.7 Simplified scheme of the peroxide curing of EPDM with dicyclopentadiene as third monomer

6.3.3.2 Co-agents in Peroxide-curing

6.3.3.2.1 General Co-agents are multi-unsaturated compounds, which are used in the peroxide-curing of elastomers. When classical co-agents, such as triallylcyanurate (TAC), trimethylolpropanetrimethacrylate (TRIM) or diallylterephthalate (DATP), are added, the crosslinking efficiency is enhanced [98-102]. Various mechanisms for the increase of the crosslinking efficiency have been proposed. In all cases a fast reaction between the

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Spectroscopy of Rubbers and Rubbery Materials co-agent and a tertiary macroradical is supposed to occur, thus suppressing unwanted side reactions such as chain scission or disproportionation. Apart from the effect on the crosslinking efficiency, the use of co-agents in peroxidecuring also imparts the molecular structure of crosslinks. It has been reported that coagents with two or more unsaturated moieties can be incorporated as individual molecules between two elastomer strands to form crosslinks [103-109]. In this way the crosslink structure of peroxide-cured elastomers can be altered. Thus, apart from the expected benefits, such as improved crosslinking efficiency, decreased compound viscosity and faster cure, the use of co-agents may also provide a tool for manipulating mechanical properties. Considering the above, it is our opinion that conclusions on the mechanism of EPDM peroxide curing in the presence of co-agents have been drawn based on a great deal of speculation. Concepts have been generally accepted, although no solid evidence for the proposed reaction pathways exists. In a number of studies the use of FT-IR spectroscopy to elucidate the mechanism of EPDM peroxide-curing in the presence of co-agents has been published. Based on spectroscopic evidence and supported by other analytical techniques, it is concluded that the actual cure mechanism [110-117] seriously diverges from the existing theories, as will be reviewed in the next paragraphs.

6.3.3.2.2 Suppression of Macroradical Side Reactions If the effect of co-agents on crosslinking efficiency is just the suppression of macroradical side reactions, such as chain scission and disproportionation, one should expect monofunctional co-agents to be as effective as their multi-functional analogues (if compared at the same molar level of unsaturation). This is definitely not the case, as will be demonstrated. In Figure 6.8 it is shown that 1-allyloxyoctane (AO) and 1,8-diallyloxyoctane (DAO) react during the peroxide-curing of EPM with the same kinetics [118]. The allyl absorptions measured directly after curing decrease to about 5% to 10% of the initial values and the ether absorptions measured after curing and extensive acetone extraction approach 100% of the initial values, indicating quantitative attachment of AO and DAO to the EPM matrix. In spite of the fact that the cure kinetics are identical, the difference compared to curing efficiency becomes apparent from the results of crosslink density measurements, as shown in Figure 6.9. The results clearly demonstrate that a diallyl functional co-agent such as DAO, provides a marked improvement in the crosslinking efficiency, whereas its monoallyl functional

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Figure 6.8 Change of IR absorptions during crosslinking of amorphous EPM at 170 °C with 2 phr of DCP and equimolar amounts of unsaturation present in 1allyloxyoctane (AO) or 1,8-diallyloxyoctane (DAO).

Figure 6.9 Effects of the addition of 1-allyloxyoctane (AO) and 1,8-diallyloxyoctane (DAO) on crosslink density, estimated from equilibrium swelling measurements in decalin

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Spectroscopy of Rubbers and Rubbery Materials equivalent, AO, does not enhance the crosslinking efficiency at all. Apparently the improvement of the crosslinking efficiency by co-agents stems from a mechanism other than the suppression of chain scission or disproportionation reactions. In efforts to come to a better understanding of the action of co-agents, the reactions of some classical co-agents in peroxide-curing of amorphous EPDM have been studied in the absence of fillers and oils, using FT-IR spectroscopy as the major tool for the elucidation of the cure mechanism.

6.3.3.2.3 Co-agent Studies In order to study the role of DATP as co-agents in the peroxide-curing of EPM, various compounds of elastomer, DCP and co-agent were cured at 170 °C in a press at different time intervals and then FT-IR spectra were recorded. Furthermore, the vulcanisates were extracted with acetone to remove the unreacted co-agent and again FT-IR spectra were recorded. Typical results are illustrated in a study using DATP as a co-agent [110]. The absorptions of the reactive allyl groups (932 cm-1) and the inert ester groups (1730 cm-1) of DATP were quantified using the CH2 vibration at 2722 cm-1 from EPM as a reference peak. In Figures 6.10 and 6.11 results from FT-IR analysis are presented. Based on the IR data of samples before extraction, it is concluded that the allyl groups react rapidly to completion within about 2 minutes, whereas the ester absorption remains constant. More allyl groups per unit of time react than peroxide radical fragments are initiated, it can be concluded that the allyl groups react predominantly via radical addition reactions, probably accompanied by radical transfer reactions. FT-IR analysis after vulcanisate extraction indicates that the co-agent is covalently bound to the elastomer matrix, as shown by the 100% recovery of the ester absorption after 2 minutes of curing. In Figure 6.10 the relative allyl absorption is divided into three parts. Part A represents the amount of co-agent attached to the elastomer matrix with one allyl group. Part B represents the amount of unreacted co-agent (the extracted co-agent fraction) and the remainder represents the amount of co-agent with which both allyl groups have reacted. Thus, a distribution of co-agent molecules which have reacted with none, one or two allyl groups, can be calculated as a function of cure time. Comparing the distribution as obtained by experiments with a theoretical distribution, in which it is assumed that all allyl groups have the same probability of reaction, a large discrepancy was established. It was found that the number of molecules which had reacted with only one allyl moiety is markedly lower as predicted according to this theoretical distribution. A similar result was found by Kloosterboer in his studies on the mechanism

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Figure 6.10 Relative allyl absorptions of EPM compounds cured with 2 phr of DCP and 3 phr of DATP, measured before and after acetone extraction [110]

Figure 6.11 Relative ester absorptions of EPM compounds cured with 2 phr of DCP and 3 phr of DATP, measured before and after acetone extraction [110]

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Spectroscopy of Rubbers and Rubbery Materials and kinetics of bulk photocrosslinking of diacrylates [119]. By using a percolation model, he explained this behaviour by showing that the reactions of diacrylates proceed inhomogeneously. Since AFM measurements showed that DATP was not homogeneously distributed in the EPM matrix (discrete spherical domains with an average droplet size of about 200 nm were found), the conclusion was drawn that the reactions of DATP during the peroxide-curing of EPM also proceed inhomogeneously. When a radical fragment is initiated within a DATP domain or in its vicinity, a local oligomerisation of allyl groups is expected to occur, resulting in the formation of a crosslinked spot within a co-agent domain, thus accounting for the relatively small fraction of allyl groups having reacted with only one allyl moiety. Since all allyl groups have reacted during curing, and in radical addition reactions (or radical transfer reactions) the active radical centre remains intact, it is proposed that the crosslinking efficiency is increased by reactions of macroradicals from the elastomer matrix with allyl groups from the co-agent domains (co-curing). This assumption was validated by a series of curing experiments in which the amount of peroxide was set at a fixed level of 2 phr and the amount of DATP was varied. A part of the cured samples was immersed in a solution of tetrahydrofuran, ethanol and 10 M NaOH at room temperature. It was demonstrated by FT-IR that after the alkali treatment, the aromatic absorptions had completely vanished. Finally, the equilibrium swelling ratio in decahydronaphthalene of the cured samples was determined both before and after alkali treatment. As expected, untreated samples showed a decrease of equilibrium swelling ratio with increasing DATP concentration. However, after alkali treatment, the equilibrium swelling ratio of cured samples was more or less independent of the amount of DATP added and the crosslink density of all samples was comparable to the sample that contained only peroxide. Thus, results are in line with the proposed mechanism that states that during peroxide-curing, bridges between co-agent domains and the EPM matrix are formed and that these bridges are the key factor enhancing the crosslinking efficiency. Moreover, it can be concluded that these bridges can be considered as extra crosslinks; the density of ‘normal’ chemical crosslinks is determined solely by the concentration of the peroxide. Apart from DATP, the curing behaviour of two isomers of DATP, i.e. diallylphthalate (DAP) and diallylisophthalate (DAIP), was also studied by FT-IR spectroscopy. Typical absorptions for the allyl and ester groups for these isomers were found at similar wave numbers, compared to DATP. Interestingly it was found that for DAP hardly any structures, comparable with that of DATP representing co-agent molecules being attached with only one allyl group could be identified. But, in spite of this fact both the compound morphology obtained rom AFM and the rate of allyl consumption as a function of cure time were similar. This behaviour was explained by assuming that

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Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy cyclopolymerisation takes place, analogous to the observations of Matsumoto and coworkers [120, 121] and Holt and Simpson [122], for the bulk polymerisation of aromatic diallylesters. Because the occurrence of cyclopolymerisation might be the reason for the different behaviour of the three isomers, attempts were made to find IR spectroscopic evidence using the following kinetic considerations: if after the addition of an EPM radical to an allyl moiety of the co-agent the next reaction was the addition of an allyl moiety of another co-agent molecule (‘normal’ polymerisation), than it is expected that at very short cure times (infinitely short) exactly one allyl group per co-agent molecule attached to the EPM backbone would be converted. However, if cyclopolymerisation were predominant, this would mean that the reaction of a co-agent molecule with an EPM radical would consume both allyl moieties instantaneously. Thus, the percentage of unreacted allyl groups per covalently bound co-agent molecule extrapolated to zero cure time provides the ratio of ‘normal’ polymerisation to cyclopolymerisation for the given compound formulation and reaction conditions. Treating the IR data accordingly, it is concluded that DAP cyclopolymerises for about 71%, DAIP for 22% and DATP for less then 5%. This result is reflected in the very low crosslinking efficiency obtained with DAP. Based on all of these observations a general scheme for the peroxide-curing of EPM in the presence of co-agents is proposed (Figure 6.12). In similar studies the kinetics and the mechanism of EPDM peroxide-curing in the presence of TAC as a co-agent has been reported [112, 113]. Peak intensities for the reactive allyl groups were monitored at 930 cm-1 and for the s-triazine ring at 822 cm-1 (out-of-plane) and 1566 cm-1 (in-plane). Again, spectra were recorded for different compounds at a fixed number of cure time intervals, both before and after acetone extraction of the cured samples. Results of these studies were in line with the previous studies on the use of diallylphthalates as co-agents. Again it was found that the peroxide crosslinking efficiency was markedly enhanced by the action of TAC, resulting from the formation of additional crosslinks across the co-agent elastomer interface. The latter observation was also supported by a study by Murgic and co-workers [113], who used a chemical probe to selectively cleave the cyanurate ester bonds after curing. Interestingly, it was found that TAC reacts via a cyclopolymerisation mechanism. It was shown that the number of reacted allyl groups per bound co-agent molecule in the initial stage of the curing process amounts to two, independent of the amount of peroxide or TAC used. Finally, mention is made to similar FT-IR studies carried out using different a,wdiallyloxyalkanes, diallylpolyethyleneglycols and methacrylic esters of

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Figure 6.12 Simplified scheme for the peroxide curing of EPM in the presence of a coagent

polyethyleneglycols [114]. Again the reactions of the reactive unsaturations were monitored by IR spectroscopy before and after acetone extraction along with the inert ether absorptions. In this case it was found that the reactions of the co-agents can be described by a ‘normal’ polymerisation mechanism. For none of these co-agents evidence for cyclopolymerisation was found.

6.3.3.2.4 Network Structure/Properties Relationships In the previous section it was demonstrated that co-agents enhance the crosslinking efficiency via co-curing of elastomer strands with co-agent domains. Because the use of co-agents alters the crosslink structure, the elastic and mechanical properties of the cured

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Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy products obtained are also affected. However, it was shown that these effects are relatively small [114]. As a general rule it can be stated that compact co-agent molecules having aromatic structures, which vitrify effectively during peroxide curing, provide worse tensile properties, because the difference in compliance between the elastomer matrix and the embedded co-agent domains is large.

6.3.3.3 Scorch Retarders in Peroxide-curing 2,6-Di-tert-butylcresol (butylated hydroxytoluene or BHT) is the most commonly applied scorch retarder in the peroxide curing of EPDM. In order to study the fate of BHT during peroxide-curing a FT-IR spectroscopic study was conducted [123]. Phenolic absorptions at 3649 and 1231 cm-1 and an aromatic absorption at 775 cm-1 were monitored as a function of cure time, both before and after extraction of the cured samples with acetone. The presence of hydroxyl and aromatic absorptions after acetone extraction indicates covalent bonding of BHT fragments to the elastomer matrix, because control experiments with uncured or sulfur vulcanised EPDM compounds showed complete extraction of BHT. The ratio of the retention of aromatic and hydroxyl absorptions after extensive acetone extraction was similar for different compound compositions cured during different time intervals. Thus, it was concluded that BHT fragments were not bound via EPDM macroradical phenoxy radical combination reactions, leaving combination via EPDM macroradicals and benzylic radicals as the most likely reaction path. In conclusion, it was stated that about 25% to 30% of the BHT became chemically bound, which was further supported by results of ageing experiments using cured samples that had been extensively extracted. FT-IR results also showed that one new (small) absorption at 1659 cm-1 appeared, which could not be attributed to peroxide decomposition products. This absorption also appeared when the peroxide-curing experiments were carried out using an amorphous EPM, indicating that the absorption did not relate to rearrangement of the third monomer moiety (ENB in this case). It is tentatively concluded that the absorption at 1659 cm-1 is related to EPDM main-chain modifications, resulting from disproportionation reactions of EPDM macroradicals with BHT radical fragments.

6.4 Concluding Remarks and Future Outlook Optical spectroscopy (IR/NMR/Raman) has been extremely useful in the study of the sulfur and peroxide crosslinking chemistry of elastomers, especially that of EPDM. The

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Spectroscopy of Rubbers and Rubbery Materials development of FT, the use of lasers and the deconvolution of spectra has enlarged the applicability of both Raman and IR spectroscopy. As a result of the high polarisability of C-S and S-S bonds, Raman spectroscopy is especially suitable for studying the sulfur vulcanisation of elastomers, whereas as a result of the high dipole moments of the polar co-agents, IR spectroscopy has been of great value in the study of the co-agent-assisted peroxide-curing of elastomers. Changes in the type of unsaturation, for instance resulting from allylic substitution by sulfur or isomerisation, have been demonstrated both with Raman and IR spectroscopy. For sulfur vulcanisation of EPDM it was shown that the relative ENB conversion (20 to 60%) is higher than often assumed. The absolute ENB conversion was shown to be governed by the vulcanisation recipe and to be independent of the EPDM type. For the ISO 4097 [82] recipe the average length of the sulfur crosslinks is 2.7 sulfur atoms. The number of converted ENB units per sulfur bridge is 2.0, indicating that crosslinks are formed predominantly. In a preliminary study it was shown that the mechanical properties of unfilled sulfur-vulcanised amorphous EPDM are determined by the chemical crosslink density. Clearly, these studies should be extended to other vulcanisation recipes and completely formulated compounds. Vulcanisation kinetics should be studied, preferably at different temperatures. It was shown that simple spectroscopic studies, which can be performed in any laboratory and not requiring high spectroscopic skills, can provide valuable data that help to understand the peroxide-curing behaviour of elastomers. Based on FT-IR studies and supported by other analytical evidence it has been demonstrated that co-agents increase the crosslinking efficiency in EPDM peroxide-curing via the formation of chemical links between the elastomer matrix and co-agent domains. These chemical links should be regarded as additional (multifunctional) crosslinks that are formed next to ‘normal’ carbon-carbon crosslinks, produced via combination of macroradicals. The latter is influenced only by the amount of peroxide used. Suppression of macroradical side reactions by co-agents, such as chain scission or disproportionation, hardly plays a role. The effect of co-agents in the peroxide-curing of EPDM is very similar to the effect of third monomers. It was concluded that the pendent unsaturation of the third monomer acts as a co-agent, i.e., the amount of third monomer governs the amount of chemical crosslinks formed by macroradical addition reactions via the unsaturated moiety of the third monomer, whereas the amount of peroxide governs the amount of crosslinks formed by macroradical combination reactions. For IR and Raman spectroscopy studies there is a clear limitation when large amounts of carbon black are applied, as is common practice in technical rubber goods. In that sense the applicability of Raman and IR spectroscopy seems to be surpassed by solid-state 13C

238

Crosslinking of EPDM and Polydiene Rubbers Studied by Optical Spectroscopy NMR spectroscopy and solid-state 1H and 13C NMR relaxation-time experiments. However, the sensitivity of solid-state 13C NMR is not as high as that of Raman and IR spectroscopy. For instance, solid-state 13C NMR of sulfur-vulcanised EPDM could only be performed when the ENB unsaturation of EPDM was fully isotopically enriched with 13 C NMR [124]. The ultimate aim of optical spectroscopy or similar studies is to elucidate the chemistry of elastomer crosslinking with respect to both the structures formed and the reaction kinetics in order to provide a sound basis for structure/property relationships. As a result, it will be possible in the future to develop rubber products with less trial and error [18, 19, 55, 56, 125].

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7

NMR Imaging of Elastomers Bernhard Blümich and Dan E. Demco

NMR imaging finds most of its applications in medical diagnostics of humans and pathology studies of animal models [1, 2]. The success of the method is based on the noninvasiveness of nuclear magnetic resonance [3–5] and the unsurpassed soft-matter contrast, which is hard to achieve with competitive methods like x-ray or computer tomography. The same advantages can be exploited in imaging studies related to materials science [6–10]. Here an important class of soft-matter materials is given by synthetic polymers above the glass transition temperature. Apart from semi-crystalline polymers like poly(ethylene), poly(propylene), some poly(amides), and polymer melts, elastomers constitute the most striking class of synthetic soft matter with a modulus similar to many biological tissues [11]. Applications of imaging for which non-destructiveness or contrast are essential, are competitive with other imaging techniques in the information gained and the cost of the experiment. Such applications to elastomers concern distributions in temperature, stress, cross-link density, modulus, and the dynamics of fluid absorption and swelling. The nucleus imaged most often is the proton. The reason is not only sensitivity but also the weak dipolar couplings between protons in a chemical group and between different chemical groups which dominate the signal decay by relaxation. These dipolar couplings are motionally averaged by the often fast but nearly always anisotropic motion of intercross-link chains. This motion and consequently the value of the residual dipolar couplings is affected by chain stiffness, cross-link density, chain orientation, temperature, additives etc. Given that the residual dipolar interactions are too strong to obtain chemicalshift resolution without sample spinning or multi-pulse techniques, relaxation techniques which probe different time regimes of molecular motion provide the primary access to contrast in imaging of elastomers. Relaxation can be probed in inhomogeneous fields, so that the homogeneous polarization field B0 is not a necessity for successful applications of soft-matter imaging. Based on this fact, small portable NMR sensors can be built, which provide NMR data from a single volume element with the same specificity as the contrast in an NMR image [12]. Such

247

Spectroscopy of Rubbers and Rubbery Materials sensors are suitable for investigations of very large objects as well as for process and quality control in an industrial setting. Their price is an order of magnitude lower than that of an NMR imager, and the spatial resolution may be several millimeters compared to about 0.1 mm in typical applications of NMR imaging to elastomer materials. The following text is divided into three parts. The first part gives a general introduction to NMR imaging and contrast. The second part reviews illustrative examples of NMR imaging to elastomer materials. The last part shows some examples of investigations of elastomer materials with mobile NMR equipment.

7.1 NMR Imaging and Contrast The power of modern NMR methods derives from the fact, that the phase of the transverse magnetisation can be measured. By use of the Fourier transformation the phase information can be converted into probability densities of resonance frequencies (Larmor Frequencies) (multi-dimensional spectra) [13, 14], densities of position (NMR images) [1–4], and probability densities of parameters like velocity and acceleration which quantifiy translational motion [15–17]. The signal phase is the angle ϕ which the precessing magnetisation M forms with the x axis of the rotating frame (Figure 7.1). In general, this angle depends on time t, because the magnitude of the magnetic field can be changed during the experiment and parts of the object can be in motion. The former is the case when pulsed field gradients (PFG) are applied to the object during measurement; the latter is the case, for example, when a polymer melt passes through an extruder or a static mixer.

7.1.1 Principle of Fourier NMR If ωL(t) is the time-dependent Larmor frequency, the precession phase (θ in Figure 7.1) can be expressed in an inhomogeneous magnetic field according to:

ϕ ( x,t ) =

t

∫

ω L ( x,t' ) dt' = − γ

0

⎡ = − γ ⎢ B0 t + ⎢ ⎣

248

t

∫ B ( x,t') dt' z

0

t

∫

0

∂Bz ( x,t' ) ∂x

⎤ x( t' ) dt' + ... ⎥ = − γ ⎥ x= 0 ⎦

⎡ ⎢ B0 t + ⎢⎣

t

∫

0

⎤ Gx ( t' ) x( t' ) dt' + ... ⎥ ⎥⎦

(1)

NMR Imaging of Elastomers

Figure 7.1 Most modern NMR techniques are based on the fact, that the phase ϕ of the precessing transverse magnetisation M(t) kann be measured. By use of the Fourier transformation the phase provides access to NMR spectra, images, and parameters of translational motion like velocity v and acceleration a. Spectroscopic parameters as well as components of translational velocity and acceleration can be used for generating contrast in NMR imaging. In the drawing the magnetisation M(t) has been generated from Mz by use of a 90° pulse of the B1 radio-frequency (rf) field in y direction

Here the inhomogeneous polarizing magnetic field points in z direction, and the spatial inhomogeneity in x direction has been expanded into a Taylor series and truncated after the second term. The first term of this series is the homogeneous field B0, the second one is the field gradient Gx = (∂Bz/∂x)x=0 in x direction. The quantity x(t) denotes the position of a magnetisation component on the x axis. For flowing and moving objects the position is a function of time. For short times also x(t) is expanded into a Taylor series and arbitrarily truncated after the third term, x( t ) = x0 + v x0 t +

1 a x0 t 2 + ... 2

(2)

The expansion coefficients are initial position x0, initial velocity vx0, and initial acceleration ax0 in x direction. Combination of equations (1) and (2) yields:

249


ϕ( t ) = ω 0 t

ϕ ( t ) = − γ B0 t t

−γ

+ kx ( t ) x0

∫ G ( t') dt' x

0

x

0 t

−γ

∫ G ( t') t' dt' v

x0

∫ G ( t') t'

1 a x0 2

x

0 t

−γ

x

2

dt'

0

− .....

+ q x ( t ) v x0

(3)

+ ε x ( t ) a x0 + .....

The individual terms of this expansion refer to different forms of ?????. The first term (γB0t) of Equation (3) concerns homogeneous magnetic fields and thus addresses NMR spectroscopy, for example, the phase contribution from the chemical shift. The second term describes the dependence of the signal phase on initial position in the presence of magnetic field gradients. It forms the basis of NMR imaging. The higher order terms describe the dependence of the phase on the transport parameters vx, ax, etc. In imaging the first, the third, and higher order terms are exploited to define contrast in NMR images. The experimental variables are the magnetic field gradient G and radio-frequency (rf) pulses, whereby 180° pulse effects a sign change of the inhomogeneous magnetic field in the rotating frame. In NMR spectroscopy the gradient Gx(t) is zero so that the signal phase is given by ω0 t. In pulsed NMR the free induction decay (FIND) s(t) is acquired as a function of time, and by Fourier transformation of the FIND, the NMR spectrum S(ω), is obtained. Similarly, to obtain an NMR image S(x0) the NMR signal s(kx) must be measured as a function of kx, the Fourier-conjugate variable of x. By the same principle, distributions S(vx0) of velocity and distributions S(ax0) of acceleration can be obtained by measuring the signals s(qx) and s(εx) and subsequent Fourier transformation. According to Equation (3) kx, qx, and εx are defined by the moments mn(t) of the gradient modulation function Gx(t), mn ( t ) =

t

∫ G ( t') t' x

n

dt'

(4)

0

Thus by suitable modulation of Gx(t) the signal phase can be made dependent on position, velocity, or acceleration. The effect of a distribution in frequency is handled either by

250

NMR Imaging of Elastomers sufficiently large gradients to dominate the phase evolution or by indirect detection schemes, where only the gradient amplitude and not the duration t of the gradient modulation period is varied (see Section 7.1.2).

7.1.2 Spatial Resolution In the absense of object motion and in a time-invariant field gradient Gx, Equation (3) simplifies to: t

ϕ ( t ) = ω L t = − γ B0 t − γ ∫ Gx (t' ) dt' x0 = ω 0 t − γ Gx t x0 = ω 0 t + kx (t ) x0 (5) 0

so that the frequency offset in the rotating frame is determined by the gradient only,

Figure 7.2 Generation of spatial resolution in NMR: (a) Conventional NMR imaging with magnetic field gradients. A magnetic field gradient Gx in x0 direction (top) converts the NMR spectrum (bottom) of an object (middle) into a projection of the object, (b) Localization of NMR signals from a large object by use of a surface coil

251


Ω = ω L − ω 0 = − γ G x x0

(6)

As a consequence, the NMR spectrum of a narrow resonance provides a projection of the object (Figure 7.2a). Each frequency Ω encodes another space position x0, and the signal intensity is proportional to the number of nuclei at this coordinate. This number is obtained by integration of the distribution of longitudinal magnetisation Mz(x, y, z) over the other two space coordinates, where the magnetic field Bz is homogeneous. Such an integral defines a projection in the mathematical sense, similar to a projection which is measured by x-ray imaging. From many such projections measured with different gradient orientations relative to the object an image of the object can be reconstructed. For reasons

Figure 7.3 Sampling principles in 2D k space: (a) Cylindrical coordinates. The angle of the field-gradient direction with respect to the x axis is given by θ = arctan{Gy / Gx}, (b) Cartesian coordinates. For rectangular gradient pulse shapes ky = -g Gy t1 and kx = -g Gx t2. Such sampling schemes are applicable to a slice which can be selected when the rf pulse is applied selectively in the presence of a gradient Gz. The areas of k space accessible by the pulse sequences shown are shaded in gray. TX: transmitter signal; RX: receiver signal; Gx, Gy: gradient signals

252

NMR Imaging of Elastomers that the FIND corresponding to the projections are measured in Fourier space, the image data are sampled in in cylindrical k-space coordinates (Figure 7.3a), and the image construction involves a conversion of the sampling grid to Cartesian k space and 2D Fourier-transformation. For a one-dimensional (1D) image of a static object like the projection in Figure 7.2a, bottom, the signal measured in a frame rotating at frequency ω0, is given by:

(

) ∫∫∫ M ( x , y , z )dz

s kx ( t ) =

z

0

0

0

0

{

}

dy0 exp i kx ( t ) x0 dx0

(7)

where kx(t) x0 = -γ Gx t x0 is the phase from Equation (5). Relaxation is neglected in this equation. Fourier transformation over kx yields the 1D image S( x0 ) =

∫∫

Mz ( x0 , y0 , z0 ) dy0 dz0

(8)

Several trajectories for k-space sampling are known in imaging, most notably sampling in Cartesian coordinates (Figure 7.3b) [1 – 5]. In this case two types of space encoding are distinguished, phase-encoding and frequency encoding. Phase encoding is only applicable to indirect detection referring to the t1 period. Here ky (cf. Equation 3, line 2) is varied in repeating scans by changing the area under the gradient pulse at constant pulse width. This procedure is more favorable than varying the gradient pulse length, because it avoids changes in the relaxation weight during the gradient pulse and changes in the signal phase from the evolution of the chemical shift or other spin interactions (cf. Equation 3, first term). Frequency encoding is used for direct signal detection during t2. Here the gradient amplitude is kept constant and as the time increases, kx increases, and data are acquired. In this case the spatial resolution is limited by relaxation, i.e. the linewidth and the width of the NMR spectrum, while for phase encoding, the spatial resolution is limited only by the signal-to-noise ratio.

7.1.3 Contrast Most conventional imaging techniques are variants of Fourier imaging in Cartesian k space. The scheme in Figure 7.3b can be modified to enable sampling of positive and negative halfs of k space in the sampling direction kx, if signal echoes are introduced in terms of gradient echoes or Hahn echoes (Figure 7.4). Hahn echoes are generated by a sequence of a 90° and a 180° rf pulse (Figure 7.4a). In the echo maximum (Rx) the effects of background magnetic field gradients and chemical shift in the phase-encoded dimension are eliminated, while for gradient echoes consisting of only one pulse

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Figure 7.4 Pulse sequences used for imaging of elastomers. A slice in direction z is defined by a frequency-selective pulse applied in a z gradient. Positive and negative halves of k space are accessed in x direction by sampling the entire echo. In the y direction the gradients are stepped through positive and negative values. The signal echo appears at the echo time tE after the first pulse: (a) Hahn-echo imaging. Signal dephasing in background gradients is refocused by application of a 180° pulse, (b) Gradient-echo imaging. Fast repetition of the experiment is enabled by small flip-angle excitation pulses α

(Figure 7.4b) they are not. Background gradients may arise from inhomogeneities of the polarizing magnetic field B0 and from susceptibility differences at interfaces within the sample. In the latter case the very sample heterogeneities to be investigated may the source for magnetic field distortions, so that gradient-echo imaging produces increased image contrast.

• Susceptibility Contrast The differences in image contrast resulting from Hahn-echo and gradient echo imaging are illustrated in Figure 7.5 by 2D images from six unvulcanized EPDM sheets separated by PTFE layers [18]. While some sample heterogeneities are recognizable in the Hahn-

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Figure 7.5 [15] Susceptibility contrast in EPDM samples at 363 K: (a) 2D Hahn-echo image acquired with the sequence of Figure 7.4a (b) 2D gradient-echo image acquired with the sequence of Figure 7.4b

echo image (a) the effects from processing by mixing and folding of layers are much better visualized in the gradient echo image (b).

• Relaxation-Time Contrast Although the spatial resolution is rarely better than 0.1 mm in NMR images of elastomer materials, NMR imaging nevertheless is useful for their analysis, because features invisible by other imaging techniques can be detected. Compared to other techniques, the number of contrast parameters in NMR imaging is abundant. Most parameters accessible by conventional NMR in homogeneous magnetic fields can be used for generation of contrast [19–21]. However, the important contrast parameters are the relaxation times, in particular those, which are sensitive to slow molecular motion. Examples are the transverse relaxation time T2, the longitudinal relaxation time T1ρ in the rotating frame, and the relaxation time T2e of the multi-solid echo decay. For example, when the space encoding sequence is preceded by a T2 filter consisting of a Hahn-echo pulse sequence, the magnetisation available for space encoding is given by:

255


⎧ t ⎫ Mz (r) = exp ⎨− E ⎬ M0 (r) ⎩ T2 (r) ⎭

(9)

where M0 is the thermodynamic-equilibrium magnetisation. Mz(r) is then used for space encoding following the principles outlined above. In Hahn-echo imaging, the T2 filter is usually integrated into the space-encoding sequence (Figure 7.4a). Parameter images are of interest in so far, as NMR parameters can be correlated with material properties either by experimental calibration or by theoretical models. An example for experimental parameter calibration is given in Figure 7.6 with a quantitative

Figure 7.6 NMR parameter image of a strained poly(dimethylsiloxane) rubber band with a cut and calibration curves: (a) Experimental curve for T2 versus strain, (b) Experimental stress-strain relationship, (c) Calibration curve for T2 versus strain obtained from combination of curves a and b, (d) Stress image obtained by recalibration of a T2 parameter image. The stress contours range from 0 to 2.4 MPa

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NMR Imaging of Elastomers stress image of a strained poly(dimethylsiloxane) (PDMS) band with a cut [22]. T2 has been determined from the tail of the signal decay curve measured with Hahn echoes for each picture cell or pixel and subsequently recalibrated with the help of the T2-versusstrain curve and the stress-strain curve determined for the rubber band without spatial resolution. The contrast variation is attributed to the strain distribution associated with the cut and to the variance of the random distribution of active silicate filler which also affects the chain mobility and thus T2. This simple calibration procedure neglects the tensorial properties of stress but nevertheless results in an image which depicts basic features of the stress distribution.

• Models of Transverse Relaxation in Crosslinked Elastomers The transverse magnetisation decay and the longitudinal decay in the rotating frame have been modelled in terms of the network topology [23–25]. The motion of chain segments in inter-crosslink chains is usually fast but anisotropic. The restricted motion leads to a non-exponential decay of the Hahn-echo maxima with solid-like contributions from residual dipolar couplings in the chain. The magnitude of these couplings increases with increasing cross-link density, decreasing temperature, and with other restrictions in

Figure 7.7 Fast but anisotropic segmental motion results in a solid-like contribution to the NMR signal. This contribution is expressed in terms of a fractional contribution q of the second moment M2 of the rigid lattice line of a single chain or residual dipolar interactions between protons. The line splitting caused by the dipole-dipole interaction depends on the orientation angle q of the internuclear vector of the coupling protons in the magnetic field B0. The distribution of orientation angles changes with the network deformation

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Spectroscopy of Rubbers and Rubbery Materials chain mobility like strain and network deformation. In the latter the case transverse magnetisation decay becomes macroscopically anisotropic, because the line splitting associated with the dipole-dipole interaction depends on the orientation angle θ of the internuclear vectors with respect to the direction of the applied magnetic field B0. (Figure 7.7). In the former case the internuclear vector of the coupling protons are isotropically distributed and no orientation dependence of the NMR signal is observed. In the Gotlib-Fedotov-Schneider model [24,26–28] the solid-like contributions are expressed in terms of a Gaussian with a fractional contribution q of the second moment M2 of the rigid lattice line of a single chain. Further signal contributions from dangling chains with approximately isotropic motional averaging are taken care of by an additional exponential relaxation term. The simplified expression suitable for analysis of imaging data is given by [24, 29] (and references therein):

⎧ t ⎫ ⎧ t ⎫ Mx (tE ) = A exp ⎨− E − q M2 tE2 ⎬ + B exp ⎨− E ⎬ ⎩ T2 ⎭ ⎩ T2 ⎭

(10)

From a fit of Equation (10) to spatially resolved relaxation curves, images of the parameters A, B, T2, q M2 have been obtained [3- – 32]. Here A/(A + B) can be interpreted as the concentration of cross-links and B/(A + B) as the concentration of dangling chains. In addition to A/(A + B) also q M2 is related to the cross-link density in this model. In practice also T2 has been found to depend on cross-link density and subsequently strain, an effect which has been exploited in calibration of the image in Figure 7.6. Interestingly, carbon-black as an active filler has little effect on the relaxation times, but silicate filler has. Consequently the chemical cross-link density of carbon-black filled elastomers can be determined by NMR. The apparent insensitivity of NMR to the interaction of the network chains with carbon black filler particles is explained with paramagnetic impurities of carbon black, which lead to rapid relaxation of the NMR signal in the vicinity of the filler particles. In the Cohen-Addad-Sotta model [23, 25] the Hahn-echo decay is expressed without the assumption of dangling chains by the expression: −1 −1 ⎧⎪ ⎛ 2 i Δ tE ⎞ ⎛ i Δ tE ⎞ ⎫⎪ Mx (tE ) = Re ⎨ M0 ⎜ 1 − ⎜1 + ⎟ ⎬ 3 Ne ⎟⎠ ⎝ 3 Ne ⎠ ⎪ ⎪⎩ ⎝ ⎭

(11)

where Δ is the average dipolar interaction on a crosslink chain and Ne is the effective number of chain segments. At sufficiently high temperature above the glass transition

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NMR Imaging of Elastomers temperature Tg, Ne can be interpreted as the number of Kuhn segments in the freely jointed chain. Thus the cross-link density is proportional to Ne-1, and the signal decay curves for rubber samples with different crosslink densities collapse onto one master curve, if Mx is plotted as a function of tE/Ne [25]. Although at very long echo times tE this approach [25] is less accurate than a refined Gotlib-Fedotov-Schneider model [26 – 28] the accuracy of Equation (11) is sufficient to analyze imaging data. A signal decay modeled by Equation (11) is depicted by a line of diamonds in Figure 7.8. A biexponential fit (solid line) fails at short tE but provides excellent agreement at long tE for the characterization of T2. At short tE a Gaussian curve (broken line) provides the best fit. Here the solid-like contribution dominates the signal decay.

• Contrast from Multi-Quantum Filtering A more detailed analysis of the NMR signal from elastomer samples, addresses finer details such as the chemical structure of chain segments. In general a hierarchy of dipolar interactions between protons exists instead of a single chain-averaged dipolar interaction [34, 35]. For example, in cis-1,4-poly(isoprene) these different dipolar interactions can

Figure 7.8 The shape of the Hahn-echo decay in cross-linked elastomers is given by the line of diamonds. In the short-time limit the curve can be approximated by a Gaussian (broken line) and in the long time limit it follows an exponential (solid line) [30]

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Spectroscopy of Rubbers and Rubbery Materials be discriminated even in static samples by multi-quantum NMR. CH groups give rise to single-quantum signals only, CH2 groups give rise to single- and double-quantum signals, and CH3 groups to single-, double-, and triple-quantum signals. In multi-quantum NMR n-quantum signals arise from coherent motion of n protons which interact by the dipoledipole coupling. Protons in close contact and on chain segments undergoing slow or highly restricted motion exhibit strong dipolar couplings, the more distant protons are coupled only weakly. Thus it is not surprising, that the multi-quantum signals are mainly produced by protons within one chemical group like CH2 and CH3 while the intergroup couplings are less significant. This is illustrated in Figure 7.9a by multi-quantum filtered 1H NMR spectra of a nonspinning sample of poly(isoprene) [34, 35]. The conventional single-quantum spectrum essentially does not provide chemical-shift resolution. However, the methylene and methyl signals can be partially separated by double- and triple-quantum filtering. The doublequantum filtered signal contains contributions mainly from the CH2 groups and a smaller contribution from the CH3 groups. The number of coupling protons defines the maximum order of the multi-quantum signal or coherence order which can arise, but any subset of the coupling spins can contribute to multi-quantum signals of lower coherence orders. For example, a double-quantum signal requires two coupling protons to flip simultaneously by absorption of two rf photons or quanta (Figure 7.9b), but also each proton can flip independent of its coupling partners in a single-quantum process. Staying with the example of two coupling protons in a CH2 group, two types of doublequantum processes can be distinguished. The double-flip and the flip-flop process. The double flip is observed by way of filtering double-quantum coherence. The resultant coherence order of the flip-flop process is zero. In contrast to established zero-quantum spectroscopy [13] it involves no signal precession at all and refers to antiparallel spin alignment of the coupling protons (Figure 7.9b). Nevertheless, it is noted that for an isolated multi-quantum pair of spins the sequence can excite only dipolar encoded longitudinal magnetisation and double-quantum coherences [34]. The dipolar encoded longitudinal magnetisation apparently behaves like zero-quantum coherence. The signals from different coherence orders can be separated by suitable phase-cycling schemes [13] of the multi-quantum filter (Figure 7.9c). In such a filter multi-quantum coherence is generated from longitudinal magnetisation by a preparation pulse sequence of duration τp, allowed to evolve in a multi-quantum evolution period t1, and then converted back to longitudinal magnetisation by a mixing pulse sequence of duration τm. Usually τp = τm. For the purpose of imaging multi-quantum filtered signals, one of the imaging schemes of Figure 7.4 follows the multi-quantum filter. Depending on the strength of the dipole-dipole interaction, the multi-quantum filtered signals are generated with different intensity (Figure 7.9d). Strongly coupled protons are

260


Figure 7.9 Multi-quantum NMR: (a) Multi-quantum filtered spectra of poly(isoprene). In contrast to the single-quantum (1Q) spectrum the double- (2Q) triple-quantum (3Q) spectra exhibit relative chemical-shift resolution. The double-quantum signal mainly arises from CH2 groups. The triple-quantum signal derives from the CH3 groups, (b) Double-quantum excitation of a methylene group generates doublequantum coherence by a double-flip process, (c) Pulse sequence for excitation and detection of multi-quantum filtered signals. The excitation time is denoted by tp. Multi-quantum signals of different orders are selected on the basis of the rf pulse phases in combination with partial signal cancellation during averaging, (d) Build-up curves of multi-quantum filtered transverse 1H magnetisation of polyisoprene: dipolar encoded longitudinal magnetisation (left) and double-quantum filtered signals (right) for τp = τm. Opposite contrast is obtained at short τp

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Spectroscopy of Rubbers and Rubbery Materials observable at short tp, weakly coupled ones at long tp. Based on the different coupling strengths of protons in different chemical groups multi-quantum signals of the same order but of different origin can therefore be emphasized and attenuated by suitable variation of τp. This fact is confirmed by the different chemical shifts observed for the double- and triple-quantum filtered signals in Figure 7.9a. But the coupling strength also depends on the modulus of the material, i. e. on cross-link density, chain extension by strain, and on temperature. In the short time limit, strong multi-quantum signals are observed from chains stiffened by high cross-link density, applied strain, and low temperature. In this regard contrast obtained from multi-quantum filtering at short τp is inversely related to contrast induced from T2 weighting. Application of the zero-quantum filter provides dipolar-encoded longitudinal magnetisation. In images of elastomers acquired with such a filter the contrast is similar to T2 contrast but more sensitive to material properties. The different contrast achieved by conventional T2 weighted Hahn-echo imaging and by multi-quantum filtered imaging for dipolar-encoded longitudinal magnetisation is illustrated in Figure 7.10 by images of a section from a silicone-breast implant envelope from poly(dimethylsiloxane) which had been implanted for 3 years [36]. The T2 weighted image (a) shows little to no signs of material deterioration, whereas in the multi-quantum filtered image (b) different islands of material deterioration are readily identified. The explanted envelope was brittle and had lost elasticity compared to a new one. Double-quantum filtered images of a strained rubber band with a cut are depicted in Figure 7.11. The image (a) has been obtained by 1H NMR with weak smoothing of the experimental data. The signal-to-noise ratio is considerably lower than for dipolar-encoded longitudinal magnetisation (cf. Figure 7.10b), and the contrast is inverted. Regions with

Figure 7.10 Images of a section from a worn silicon-breast implant envelope: (a) T2weighted imaged, (b) image of dipolar encoded longitudinal magnetisation

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Figure 7.11 Double-quantum images of a strained rubber band with a cut: (a) 1H image, (b) 2H image of deuterated spy molecules incorporated into the rubber network by swelling, (c) finite element simulation of the stress distribution

weak dipolar couplings appear bright in Figure 7.10b and dark in Figure 7.11a. Thus the region of high strain in the center of the cut shows high signal intensity in the doublequantum image. Similar information can be obtained from double-quantum filtered imaging of deuterons [37]. From the point of NMR spectroscopy deuterons behave like a pair of coupled protons. Deuterated butadiene oligomers have been incorporated into a rubber band as spy molecules by swelling and subsequent solvent drying. The geometry of molecular motion of the spy molecules probes the void geometry of the rubber network, which becomes anisotropic upon straining the sample. The motionally averaged residual quadrupolar coupling of the deuteron is sensitive to the anisotropy of molecular motion in a way similar to the dipole-dipole interaction between two protons, except that it is zero in isotropic elastomers, while the residual dipolar coupling of an intercross-link chain is not. In either case multi-quantum NMR can be used to probe the residual dipolar and quadrupolar interactions which are influenced by the state of the rubber network. The stress distribution obtained for a strained rubber band by double-quantum deuteron NMR (Figure 7.11b) [37, 38] agrees with the one obtained from proton NMR (Figure 7.11a) [39]. The quality of the deuterium image appears better because of much stronger data smoothing. The measured signal intensities match the stress distribution simulated by finite element methods (Figure 7.11c). Like the stress image from T2 data (cf. Figure 7.6) the double-quantum filtered images do not account for the orientation dependence of the stress tensor. To this end multi-quantum filtered NMR images need to be measured from different orientations of the strained sample in the magnetic field [40], and the resultant data set should be analyzed for the principle components of the stress

263

Spectroscopy of Rubbers and Rubbery Materials tensor, that is, for the trace, the anisotropy, the main principle value, and the orientation angles in the sample. This approach to imaging of anisotropic material properties has been demonstrated in diffusion tensor imaging of ordered biological tissue [41 – 43].

7.2 Applications Applications of NMR imaging outside medicine are relatively rare [6, 7]. Potentially rather useful applications are in the elastomer industry. This section features some selected examples which illustrate the type of information obtainable by imaging of elastomers. Heterogeneities in technical elastomers arise in different stages of elastomer production and product use. Even in a perfectly prepared homogeneous elastomer product, unavoidable ageing processes induce space-dependent defects. Examples for sources of defects are: • Mixing process: Technical rubbers are blends of up to about 30 different compounds like natural rubber, styrene-butadiene rubber, silicate and carbon-black fillers, and mobile components like oils and waxes. These components show a large variety of physical, chemical, and NMR properties. Improper mixing leads to inhomogeneties in the final product with corresponding variations in mechanical and thermal properties (cf. Figure 7.4). • Vulcanisation: The changes in segmental mobility from progress of the vulcanisation reaction and from the associated sample temperature distribution can be monitored directly by NMR imaging in situ [44]. Heterogeneous structures arise from effects of thermal conductivity, which lead to space-dependent temperature profiles during the vulcanisation process depending on the position of the heat source and on heat dissipation. As a result inhomogeneous cross-link densities may be established [45]. In the covulcanisation of blends and sheets from different formulations inhomogenities in cross-link density may arise from differences in solubility and diffusion of the curatives [46]. Different components of rubber blends can be mapped by exploiting the editing capabilities of motionally narrowed, swollen rubber samples [47] and of proton-detected 13 C imaging, which are useful even in the case of unresolved lines [48]. • Ageing: Ageing processes are most often introduced by UV irradiation, exposure to heat and oxygen, and by biological mechanisms (cf. Figure 7.10). Depending on the load applied, different ageing processes are observed [49]. Thermal oxydative ageing usually leads to the formation of hardened surface layers in natural rubber (NR) as well as in synthetic rubber (SBR, styrene-co-butadiene rubber) [50 – 54]. Typically these layers approach a thickness of 200 to 300 μm and inhibit the progress of the ageing process further into the bulk of the sample. The material hardening is explained

264

NMR Imaging of Elastomers by an increase in cross-link density. In the absence of oxygen, chain scission may dominate at elevated temperatures with an associated increase in segmental mobility. Other types of ageing involve aggressive fluids and gases. In this context a sample of degraded rubber hose has been investigated [55], but also the degradation of polyethylene pipes [56], and the enzymatic degradation of biologically synthesized polymers [57] have been studied through NMR imaging. Related investigations have been carried out on asphalts [58]. Ageing associated with swelling of the rubber particles has been observed in crumb-rubber modified asphalts [59]. • Mechanical load: Static mechanical load by strain or compression leads to stretching of random-coil polymer chains in the direction of sample elongation and chain compression in the orthogonal directions (cf. Figure 7.7). Stress and strain effects can be analyzed for instance by parameter maps of T2 [22] (cf. Figure 7.6), and by 1H [34, 35] and 2H [37, 38] multi-quantum imaging (cf. Figure 7.11). Dynamic mechanical load leads to sample heating where the temperature distribution in dynamic equilibrium is determined by the temperature-dependent loss-modulus and the thermal conductivity of the sample. Because T2 scales with temperature for carbon black filled SBR, a T2 map provides a temperature map of the sample. Such temperature maps have been measured for carbon-black filled SBR cylinders for different filler contents and mechanical shear rates [60].

7.2.1 Defects and Heterogeneities in Technical Elastomer Products A common source of sample heterogeneities arises from filler clusters. Hardly any technical elastomers is free of it. Even the variance of the distributions statistics of active filler may lead to heterogeneities detected as an average over a volume cell (or voxel) with the spatial resolution of NMR imaging (cf. Figure 7.6d). Small filler clusters may be invisible by Hahn-echo imaging but may give rise to enlarged imaged distortions from susceptibility effects in gradient-echo imaging (cf. Figure 7.5). An example for elastomer defects is shown in Figure 7.12 by images of a rubber gasket which had been exposed to oil, heat, and pressure in an overload test. No signs of failure can be detected by visual inspection of sections cut out of the gasket (a). However, in a T2 weighted spin-echo image (b), filler defects can be identified (circular structure in the center) as well as regions of severe sample hardening (dark strip on the right), and a swollen periphery, where the technical oil has penetrated into the gasket (bright contour lining the sample). In fact closer analysis reveals a chromatographic separation of the swelling fronts associated with different components of the technical oil [61]. In the image (c) from a different section of the gasket a crack is noticeable in the middle of the image, and the strips of hardened material have grown larger.

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Spectroscopy of Rubbers and Rubbery Materials The visibility of defects may be enhanced by swelling in a suitable solvent. This is illustrated in Figure 7.13 by images of a filled and cured sample of the terpolymer isobutylene-pmethylstyrene-p-bromoethylstyrene [47]. In the image (a) acquired with an echo time tE = 2 ms a few large inhomogeneities identified as voids are recognizable. Much more defects become oberservable by swelling in cyclohexane, which softens the matrix and increases the transverse relaxation time for better signal-to-noise ratio. The position of the voids remains unchanged, so that swelling preserves essential features of the elastomer

Figure 7.12 Images of a carbon-black filled rubber gasket: (a) Photograph, (b) T2weighted image from an undamaged region, (c) T2-weighted image from a region damaged in an overload test from pressure applied in a bath from technical oil

Figure 7.13 1H Hahn-echo images of a terpolymer: (a) Unswollen sample, (b) sample after swelling in cyclohexane. The images were taken from a 15 mm diameter cylinder (adapted from [47] with permission of the authors)

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NMR Imaging of Elastomers matrix. This approach to imaging provides the opportunity to optimize contrast in a sample-specific way by imaging the unswollen polymer, the swollen network, and different solvents with chemical-shift-selective excitation. Samples with very short relaxation times are difficult to image by Hahn-echo methods. Images of such samples can be acquired by single-point imaging [62, 63], and its variant, the SPRITE (single-point ramped imaging with T1 enhancement) method [64]. Basically a small-flip-angle rf pulse is applied in the presence of a magnetic-field gradient, and a single data point is acquired a short time after the pulse. Then the experiment is repeated with a different value of the field gradient vector. In this way k space is scanned point by point. At first sight the method appears to be inefficient in terms of acquisition time. But Hahn echoes are avoided and small flip angles are used, so that repetition times are fast. Moreover, the time delay between the pulse and data acquisition can be varied to introduce chemical-shift modulation in the acquired signal. If a refocusing pulse is placed in the middle of the acquisition delay, the chemical-shift evolution is refocused and just a T2 weight is introduced. 128 x 64 point SPRITE images of elastomers can be acquired in about three minutes including signal averaging. An example of such an image is given in Figure 7.14 for a tyre section consisting of carbon-black filled SBR layers and an NR

Figure 7.14 T2* weighted image of a car-tyre section showing layers of SBR (left) and NR (right) which are separated by polymer fabric. The image was acquired by the SPRITE technique in 200 s

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Spectroscopy of Rubbers and Rubbery Materials layer separated by textile fabric. The acquisition delay of 5 ms results in a T2* weight which enhances the signal from the softer NR component [65]. Spin density images are obtained with this method for samples with long relaxation times compared to the acquisition delay. Two T2-weighted Hahn-echo image from sections of the tyre treads are compared in Figure 7.15. The samples had to be cut from the steel belt in order to avoid severe distortions of the polarizing magnetic field by the belt. In each image the tyre surface is on the left (hard material, dark) and the base is on the right (soft material, bright). Next to the base two layers of reinforcing polymer fibers (line of dots) are embedded in a hard formulation in image (a) and in a soft formulation in image (b). Furthermore, in image (b) the soft base material bulges out into the tread material, and filler clusters can be identified. The performance of tyre (b) was inferior to that of tyre (a), where the base material is hard. Overlaid to the images is a variation of T2 which had been measured by the NMR-MOUSE (see Section 7.3 and Figure 7.22) at different depths starting from the right and from the left of the sample. Extreme differences between the signals are found at 6.5 mm depth in this case, which correlate with the hard and soft materials. This example demonstrates, that the data measured by the NMR-MOUSE can deliver information which correlates with that obtainable by NMR imaging. However, in comparison with the imaging data, the signal-bearing volume is much larger and less well defined. But the NMR-MOUSE measurements are nondestructive and can be

Figure 7.15 T2 weighted images of car-tyre sections showing in each image the tread (left) and the base with polymer fibers (right): (a) Hard base, (b) Soft base. The overlaid curves are T2 data measured as a function of depth by the NMR-MOUSE (see Section 7.3.1)

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NMR Imaging of Elastomers performed on tyres with steel belts, so that they can be carried out during tyre testing, for example, on the race track.

7.2.2 Covulcanisation During fabrication of elastomer products like car tyres different elastomer formulations often need to be covulcanized. At the interface between an SBR and an NR layer, an interfacial layer with a modulus higher than either SBR and NR had been detected in a sample from a used car tyre [51]. Subsequently such interfaces were detected in other samples as well. For example, in unfilled elastomer composites they can be seen even by visual inspection [52]. In Hahn-echo images the interface is identified by a dark line paralleled by a slightly brighter line (Figure 7.16a). The contours associated with spin density and fast relaxing components are eliminated from the image by forming the ratio between two Hahn-echo images acquired with different echo times, a short and a long one. While the image in Figure 7.16a has been acquired with an echo time of tE = 3.3 ms, the image Figure 7.16b has been obtained by normalizing the image acquired with

Figure 7.16 T2 weighted Hahn-echo image: (a) and quotient image (b) computed as the quotient of two Hahn echo images acquired with different echo times of an unfilled SBR/NR covulcanisate with dimensions 9 mm x 30 mm. The acquired signal has been integrated over the sample thickness of 1 mm. Contrast in the quotient image is determined mainly by transverse relaxation. Contrast in the Hahn-echo image is formed by a mixture of spin density and relaxation

269

Spectroscopy of Rubbers and Rubbery Materials tE = 4.2 ms to an image acquired with tE = 3.4 ms. In the latter image the interface appears abrupt and well defined. The visual impression of the interface is similar to that observed in Hahn-echo images (Figure 7.17a). In order to obtain further details about the interfacial region, 1D profiles were acquired with single-point imaging and different contrast filters (Figure 7.17b) [65, 66]. If the acquisition delay inherent to single-point imaging is chosen appropriately (0.35 ms), the signal decay is affected by the weak but nevertheless effective chemical-shift modulation, so that the image contrast is dominated by chemical shift effects (right scale, smooth line). This image shows a rapid transition between both components and a narrow

Figure 7.17 Interface between covulcanized sheets from unfilled SBR and NR: (a) Photograph, (b) 1D NMR images acquired with single-point imaging and different contrast filters. Smooth line: chemical contrast (right scale). Rugged line: relaxation contrast (left scale)

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NMR Imaging of Elastomers chemical interface with a width of less than 0.1 mm. A 180° pulse in the middle of the acquisition delay eliminates the chemical shift modulation and the signal decay is defined only by transverse relaxation. The contrast in the quotient profile corresponding to an image acquired with an intermediate echo time (acquisition delay) normalized to an image acquired with a short echo time (rugged line, left scale) provides mainly relaxation contrast and thus information about molecular mobility or sample hardness. This image of the physical interface shows a broad transition from the hard SBR component to the soft NR component with a width of the order of 0.5 mm. This is the space scale on which the modulus changes. More complicated shapes of the interface can also be observed. The shape and dimension of the interface are defined by the concentration differences in the vulcanizing agents, which diffuse at elevated temperatures across the interface until their diffusion is hampered by their role in the vulcanisation reaction. Thus the interface arises from a delicate balance of diffusion, reaction, heat supply, and removal.

7.2.3 Blending Interfaces similar to those encountered during covulcanisation may arise in blends with incomplete mixing. In unvulcanized samples heterogeneities from blending can be detected by gradient echo imaging, where the contrast is enhanced by differences in magnetic susceptibility of the components (cf. Figure 7.5). However, fine structures are usually homogenized during the vulcanisation process. Larger structures survive, in particular, when one component had already been vulcanized. An example of such a case is illustrated in Figure 7.18 by two orthogonal slices through a block of a vulcanized blend from a soft (bright) and a previously vulcanized hard (dark) component. In slice (a) the structures invoked by the mixing of components appear course but random, but slice (b) reveals stream lines of material flow in the rolling mill.

7.2.4 Crosslink Density Cross-link density and parameters relating to the network structure can be measured by NMR by analysis of the transverse relaxation decay (cf. Section 1.3) and the longitudinal relaxation in the rotating frame [67]. Combined with spatial resolution, the model-based analysis of relaxation yields maps of cross-link density and related parameters [68]. Often the statistical distribution of relaxation parameters over all pixels provides a reduced data set with sufficient information for sample characterization and discrimination [68]. Relaxation curves can be measured by simple NMR experiments. On the other hand, information about cross-link density and chain stiffness can be retrieved by exploiting

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Figure 7.18 T2 weighted orthogonal slices of a cabon-black filled rubber blend with a soft (bright) and a hard (dark) component. Insufficient blending in a rolling mill is visualized by the stream lines in slice (b) residual dipolar couplings (cf. Section 1.3). Experiments to measure such couplings are more sophisticated, but better image contrast can be obtained and more specific structural information such as the dynamic order parameter of individual chemical groups within the network [35]. Model-free data of cross-link density can be obtained by NMR when calibrating NMR parameters against values of cross-link density determined by other methods. However, NMR relaxation and residual dipolar couplings are sensitive not only to cross-link density but also to restrictions in chain motion in general, so that effects from temperature variation, strain, and cross-link density need to be separated based on a priori information. Most important, comparative measurements on different samples near room temperature need to be carried out at the same temperature or extrapolated to a reference temperature based on previously determined temperature coefficients. Conventionally, cross-link density is determined by measurements of the modulus, the glass transition temperature Tg, and by solvent uptake in swelling experiments. In these procedures, the chemical cross-link density cannot be discriminated from network-

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NMR Imaging of Elastomers filler interactions. If NMR experiments are performed over 90°C above Tg, physical interactions imposed by chain entanglements exert no influence on transverse NMR relaxation. Moreover, the presence of carbon-black filler usually does not affect the NMR data in contrast to functionalized silicate filler. This is attributed to rapid signal relaxation of network chain segments near carbon-black filler particles caused by paramagnetic centers in the filler. Thus for the technologically important class of carbonblack filled elastomers, NMR can provide the chemical cross-link density.

7.2.5 Vulcanisation Process Variations in cross-link density may arise from spatial variations in the rubber formulation, although short-scale variations are often smoothed by component diffusion during the vulcanisation process. Differences on the mm scale can lead to interfacial structures like those depicted in Figures 7.16 and 7.17. Another source of variations in cross-link density on the mm scale is the curing process in combination with the sample geometry. Heat is supplied to the sample for a certain time and after vulcanisation is removed from the sample in a certain time. Near the heat source vulcanisation sets in first, and near the heat sink it sets in last. Depending on how the heat is supplied to and withdrawn from the object, complicated time-dependent temperature profiles are established in the sample. The vulcanisation process has been followed across a simple 1.4 mm thick disc from unfilled SBR by NMR imaging of the transverse signal decay [44]. The measurements were carried out in a magnetic field slightly inhomogeneous from the construction of a special vulcanisation probe, and the transverse relaxation time T2* in the inhomogeneous magnetic field was measured. Measurement of T2 would have been too time-consuming to follow the vulcanisation process. Despite accelerated signal decay from the field inhomogeneities, T2* turned out to still be sensitive to chain mobility because of a sufficiently strong contribution from T2. A somewhat nonlinear relationship between the transverse relaxation rate 1/T2* and the reduced shear modulus was found. Nevertheless, the relationship that low modulus corresponds to high T2* and high modulus to low T2* can be used to interpret the imaging data across the SBR sheet measured during the vulcanisation process (Figure 7.19). Immediately after sample heating (5 min.) chain mobility is high and so is T2*. With the formation of cross links chain mobility decreases and so does T2*. The vulcanisation front is defined by the center of change in T2*. It can be seen to migrate through the thin sample on a time scale of about half an hour with temperatures in the range between 140° and 170° C across the sample. With increasing time the slope of the T2* curve flattens, indicating a broadening of the reactive vulcanisation zone.

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Figure 7.19 Time-resolved T2* parameter images across a 1.4 mm thick sheet of SBR following the vulcanisation process

7.2.6 Ageing Ageing of elastomers is a process which affects the mobility of intercross-link chains by packing, chain scission, and formation of new cross-links. It can therefore be studied by the same methods as cross-link density. Most conveniently ageing is investigated by relaxation methods. Thermo-oxydative ageing of many elastomers including SBR and NR leads to the formation of a brittle surface layer with reduced mobility which appears dark in Hahn-Echo images. The aged region asymptotically approaches a thickness of typically 0.3 mm with time [69]. Often this region is followed by a zone of brighter image intensity, which can be attributed to chain scission or the accumulation of low molecular-weight additives. The time evolution of the thermo-oxydative ageing in SBR has been studied by NMR imaging [52]. In order to detect the aged surface layer with reduced signal intensity in Hahn-echo images, two aged sheets were stacked, and the stack was imaged along its axis. In this way the signal from the soft core delineated the hardened surface layers (Figure 7.20a). The signal loss from the hardened surface regions was specified in terms of the area A1 in the image which appeared dark relative to the total area A1 + A2 of the object in the image. This quantity defines the ageing parameter α. It was found to follow a simple exponential law with time for a sample with and a sample without ageing

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Figure 7.20 Analysis of thermal oxydative ageing by T2-weighted imaging (a) Two SBR layers stacked after ageing. The aged layer suffers a signal loss marked by the area A1. The relative signal decrease defines the ageing parameter a. The progress in ageing can be followed by evaluating the ageing parameter as a function of ageing time

protectant (Figure 7.20b). For both samples the time constant is the same, so that the ageing reaction proceeds with the same speed. However, the amplitudes are different, so that the number of chain segments affected by the ageing reaction is different. Less segments are exposed to the oxydative ageing attack in the sample with the ageing protectant, confirming the efficiency of the ageing protectant as a radical scavenger.

7.2.7 Sample Deformation Sample deformations modify the number of accessible conformations of intercross-link chains (cf. Figure 7.7), so that they can be detected by analysis of relaxation and residual dipolar couplings. This is illustrated for strained rubber bands with a cut in Figures 7.6 and 7.11. Dynamic mechanical load on elastomer products is often exerted at small deformations and low deformation rates but over extended time periods. Then part of the mechanical energy is dissipated into heat depending on the value of the loss modulus. As a consequence, a temperature profile is established within the sample. Then the modulus

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Spectroscopy of Rubbers and Rubbery Materials varies across the sample depending on the temperature profile, and properties determined for thick samples under dynamic load are averaged quantities. The temperature profile associated with sample heating during weak dynamic shear deformation of carbon-black filled SBR cylinders of 10 mm in diameter and 10 mm in height has been imaged by NMR by use of a specially designed probe [70]. The transverse relaxation time strongly depends on temperature (Figure 7.21a), so that temperature can be mapped by parameter imaging of T2. Axial parameter projections have been acquired in dynamic equilibrium at a shear rate of 10 Hz and a pixel resolution of 0.4 x 0.4 mm2 for carbon-black contents ranging from 10 to 70 phr. One-dimensional cross-sections through those projections are depicted in Figure 7.21b. An increase of the temperature in the center of the sample is observed with increasing carbon-black contents which scales with the increasing loss modulus of the samples. The 70 phr sample is warmer by over 10° C in the center of the sample than the 10 phr sample.

7.3 Spatially Resolved NMR In a factory environment NMR instrumentation needs to be robust and possibly mobile for quality and process control at different stages of product development, fabrication,

Figure 7.21 Temperature profiles from T2 parameter images of SBR cylinders with different carbon-black filler contents undergoing oscillatory shear deformation: (a) Temperature calibration curves, (b) Temperature profiles across the cylinders

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NMR Imaging of Elastomers and quality control. These demands are difficult to fulfill with sophisticated pulse sequences and highly homogeneous magnetic fields B0. For this reason low-resolution NMR is well established in industrial laboratories, for example in the food, the cosmetics, and the polymer industry [71]. In low-resolution NMR signal amplitudes, relaxation times, and diffusion constants are measured. These parameters are most important in definition of contrast in imaging of elastomers, but can be measured also in inhomogeneous magnetic fields. Therefore, permanent magnets can be employed at lower field strength, where transverse relaxation is even more sensitive to slow molecular motion than at high field. Low-resolution NMR in inhomogeneous magnetic fields has been pioneered in the oil industry for well logging by single-sided inspection of rock formations [72, 73]. The same principles of single-sided NMR [74] are applied by the NMR-MOUSE, a mobile universal surface explorer which weighs one to three kilograms with which NMR relaxation parameters can be acquired non-destructively from surface-near volume elements of arbitrarily large objects [75 – 80]. The NMR-MOUSE provides interesting applications for quality assessment of elastomer products. Because of the simplicity of the device and the pulse sequences, it is suitable for use in a factory and can be transported to the object of investigation for spatially resolved NMR of accessible sample regions.

7.3.1 The NMR-MOUSE The NMR-MOUSE is a portable NMR sensor which works in highly inhomogeneous magnetic fields. Because of field inhomogeneity NMR spectroscopy of the chemical shift is not readily possible, but relaxation times and parameters of translational motion can be measured by echo techniques. These are the most important NMR parameters which are exploited for contrast in imaging. Unless fluids are investigated field inhomogeneities are essentially no obstacle for relaxation analysis [80], because molecular motion by translational diffusion is absent. Basically the NMR-MOUSE consists of an u-shaped magnet with an rf coil in the magnet gap (Figure 7.22a). In a portion of the accessible space above the gap, the field lines of the polarizing magnetic field B0 and the rf field B1 possess orthogonal components. In this region of space the NMR effect can be invoked. The device can be built rather small (Figure 7.22b) and the field profile is not severely affected by magnetic components in the vicinity of the device. For example, NMR measurements can be executed on intact tyres with steel belts and on conveyor belts with steel cords (cf. Figure 7.25). Even the laboratory weathering of a 0.5 mm thick poly(vinyl chloride) coating on a 1.5 mm thick sheet of iron could be followed with the NMR-MOUSE [76]. The polarizing magnetic field possesses field gradients of the order of 10 to 50 T/m [78]. In such large field gradients, even a short rectangular pulse is a selective pulse and no free induction decay signal can be observed with a minimum instrument deadtime of about 10 μs after an rf

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Figure 7.22 The NMR-MOUSE: (a) Schematic. The NMR sensor consists of an u-shaped permanent magnet with a solenoidal rf coil placed in the gap. (b) Photo of the NMRMOUSE testing a tyre. (c) Example of a train of successive Hahn echoes generated according to Carr, Purcell, Meiboom and Gill (CPMG echo train) for carbon-black filled SBR measured by the NMR-MOUSE. The time constant of the echo-envelope defines T 2

pulse. Therefore echo techniques need to be employed, which eliminate the phase differences from signal precession in different magnetic field strengths. Such echo techniques are the Hahn echo and its multi-pulse version, the CPMG sequence. Other techniques are the solid echo and its multi-pulse version, the OW4 sequence [80] (and references therein). Figure 7.22c gives an example of a CPMG echo train from carbonblack filled SBR. In theory, the echo maxima follow a decay, which has been discussed in Section 1.3. In practice, the particular pulse sequence has an influence on the shape of the echo decay function, because the signals are acquired in highly inhomogeneous 278

NMR Imaging of Elastomers magnetic fields [81], and relaxation curves measured by the NMR-MOUSE are often evaluated in terms of a mono- or biexponential decay. A mono-exponential decay function is schematically indicated in Figure 7.22c. The time constant of this decay is the effective transverse relaxation time T2. It can be calibrated against material properties.

7.3.2 Applications Some selected applications of the NMR-MOUSE to elastomers are summarized in Figure 7.23. In (a) T2 has been measured for a series of carbon-black filled NR samples

Figure 7.23 Applications of the NMR-MOUSE to elastomer materials: (a) T2 values for a curing series of carbon-black filled NR. Comparison of values obtained at high homogeneous field (DMX 300) and with the NMR-MOUSE, (b) cross-link series of unfilled SBR with different sulfur content, (c) T2 versus glass-transition temperature Tg of unfilled SBR by the CPMG and the steady-state CPMG methods, (d) normalized Hahn-echo decay curves for poly(butadiene) latex samples. Different decay rates are obtained for small medium and large cross-link densities

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Spectroscopy of Rubbers and Rubbery Materials with different curing times. Normalized values measured at high and homogeneous magnetic field (7 T, DMX 300 NMR Spectrometer) are compared to those measured with the NMR-MOUSE (0.5 T). Although the values differ because of different B0 field strengths, they closely follow the same trend. This confirms, that relaxation measurements by the NMR-MOUSE are a valid alternative to relaxation measurements at homogeneous magnetic fields. In (b) relaxation times are shown for a series unfilled SBR samples with variations in sulfur content. With increasing sulfur content the chemical cross-link density increases. Relaxation times are given for measurements with Hahn echoes and with solid echoes. The change in relaxation times and thus the sensitivity of the method is greatest for the Hahn-echo measurements at small cross-link density and for the solid-echo measurements at high cross-link density. Solid echoes reduce the signal attenuation from the dipoledipole interaction between two spins in addition to producing a Hahn echo. At high cross-link density multi-center dipolar couplings become effective, and the solid echo becomes less effective in reducing signal attenuation from dipolar interactions. The Hahn echo does not affect the dipole-dipole interaction at all. This explains the difference in contrast obtained with both methods. A reduction in measurement time is gained, when the measurements are performed in dynamic equilibrium between rf excitation and longitudinal relaxation [13]. NMRmethods which operate in this regime are referred to as steady-state methods. They deliver equivalent information compared to methods which demand complete relaxation to thermo-dynamic equilibrium between scans [13]. T2 measurements performed with the CPMG and the steady-state CPMG method are depicted in Figure 7.23 (c) and correlated with the glass transition temperature Tg of unfilled SBR laboratory samples of different cross-link density. The NMR measurements were done at room temperature, whereas Tg had been determined by temperature-dependent measurements of the dynamic-mechanical loss modulus [75,76]. The correlation between both measurements is not surprising, because both probe the network dynamics. These data demonstrate, that in certain cases, Tg can be determined locally and nondestructively on large samples at room temperature by the NMR-MOUSE. The large magnetic-field gradients give rise to rapid signal loss from molecules with translational motion, so that signals from low molecular-weight fluids are suppressed at echo times tE of the order of 1 ms and more. The signal detected at larger echo times is from larger molecules or particles. This effect of solvent suppression is exploited in the characterization of cross-link density in poly(butadiene) latex samples (see Figure 7.23 (d)). The signal decay for weak cross-link density is slow compared to that for large cross-link density. In summary, Figure 7.23 demonstrates, that the NMR-MOUSE is a suitable device for determination of relative cross-link density in a number of different soft materials.

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NMR Imaging of Elastomers In elastomer samples with macroscopic segmental orientation, the residual dipolar couplings are oriented as well, so that also the transverse relaxation decay depends on orientation. Therefore, the relaxation rate 1/T2 of a strained rubber band exhibits an orientation dependence, which is characteristic of the orientational distribution function of the residual dipolar interactions in the network. For perfect order the orientation dependence is determined by the square of the second Legendre polynomial [14]. Nearly perfect molecular order has been observed in porcine tendon by the orientation dependence of 1/T2 [77]. It can be concluded, that the NMR-MOUSE appears suitable to discriminate effects of macroscopic molecular order from effects of temperature and cross-link density by the orientation dependence of T2. The NMR-MOUSE is simple mobile NMR sensor suitable for operation in an industrial environment. Its use for quality control is demonstrated in Figure 7.24 by the T2 statistics measured at both sides of a conveyor belt with steel cords. The B0 distortions from the steel cords were minimized by suitable orientation of the NMR-MOUSE with respect to the direction of the cords. The upper side of the belt exhibits a higher average T2 value than the lower side, which indicates lower cross-link density. In addition to that, the variance of the measured values is larger for the lower side, indicating a better quality material on the upper side. Depending on the size of the rf coil and the size of the sensitive volume, typical measurement times at room temperature range from ten seconds to some ten minutes. Following suitable instrument optimization and adaptation to specific processes, the NMR-MOUSE promises to be of use also in process control.

Figure 7.24 Quality control of a conveyor-belt section with steel cords: (a) Position of measurement points. (b) T2 values for upper and lower sides determined with the CPMG sequence

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7.3.3 Imaging with the NMR-MOUSE Images can be obtained with the NMR-MOUSE by measuring voxels individually by lateral displacement of the device or a change of excitation frequency to shift the sensitive volume in depth (cf. Figure 7.15). On the other hand, also pulsed magnetic field gradients can be employed for phase encoding of the space information. Frequency encoding is hampered by the nonlinear magnetic field profile. In direction parallel to the magnet gap, the magnetic field gradient is weakest, and a sensitive volume of a centimeter and more in width can be excited by a short rf pulse. Solenoidal gradient coils can readily be incorporated into the gap and pulsed to produce antiparallel magnetic fields (Figure 7.25b). Then a gradient field is established along the gap, and spatial resolution can be introduced into the measurement by phase-encoding techniques similar to single-point imaging. A suitable pulse sequence is shown in Figure 7.25a. The effects of the background field inhomogeneity Bz – B0 and of chemical shift on the magnetisation phase are balanced in the peak of the Hahn echo, and the only phase evolution from the pulsed field gradient remains. By applying gradient pulses with different amplitude, k space can be scanned in the direction along the gap, so that Fourier transformation of the acquired signal produces a 1D image [66]. This measurement protocol has been applied in a study of a rubber sheet with parallel textile fibers (Figure 7.25c). The fibers do not contribute to the detected signal, so that their positions can be located by the dips in the 1D profile. Clearly, the basic imaging scheme can be extended to include relaxation-time contrast for discrimination of variations in cross-link density and strain, and the 1D MRI-MOUSE (magnetic resonance imaging MOUSE) can be extended with further gradient coils to permit imaging in three dimensions. Numerous applications of the MRI-MOUSE can be envisioned in soft matter analysis, in particular in those areas, where imaging with conventional equipment has proven to be successful, and where smaller, less expensive, and mobile devices are in need.

7.4 Summary NMR imaging has been tested in a multitude of cases and many simple and sophisticated methods have been developed. The most successful applications of the method, however, are in the field of soft matter, because transverse relaxation times are compatible with current hardware technology. In the non-medical field, NMR imaging of elastomers shows great promise for genuine applications, and in addition to academic laboratories, the method is used in industry for product development and control. While academic efforts often strive towards the development more sophisticated measurement schemes and

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Figure 7.25 One-dimensional imaging with the NMR-MOUSE: (a) Single-point imaging sequence for phase-encoding of space, (b) drawing of the NMR-MOUSE with coils for pulsed field gradients, (c) sample of an elastomer sheet with parallel textile fibres and one-dimensional NMR image with the space direction perpendicular to the fiber direction

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Spectroscopy of Rubbers and Rubbery Materials hardware tools, the demands for routine industrial use are simplicity, precision, and economy. In all three demands, conventional NMR usually does not score high: Most modern NMR methods require many precise pulses and highly homogeneous magnetic fields. The precision of the method is often limited by inherently low sensitivity of NMR. Economy often sets preferences against NMR, because the method requires expensive equipment and personnel. Nevertheless, for elastomers simple methods like Hahn-echoes and gradient echoes are useful for materials characterization and imaging. Homogeneous magnetic fields are not a prerequisite for imaging and relaxation measurements, and inexpensive devices like mobile low-field spectrometers, the NMR-MOUSE, and mobile imagers can be built for use near or in the production line and for operation by technicians. For these reasons, NMR of elastomers including NMR imaging of elastomers appears to address one of the industrially most relevant applications to NMR. Further developments of NMR instruments and the growing understanding of NMR in inhomogeneous fields are expected to be of beneficial impact for increased use of the method in engineering and fabrication.

Acknowledgements Continuous support of this work by Deutsche Forschungsgemeinschaft (DFG) and Fonds der Chemischen Industrie (FCI) is gratefully acknowledged.

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NMR in Soft Polymeric Matter: Nanometer Scale Probe Jean Pierre Cohen Addad

8.1 Introduction Thinking of plastics all around us, polymers are usually regarded as materials substituted for wood, glass or metal even though macromolecules also go into soft matter such as rubber, latexes, contact-lenses, thin films; polymeric substances may be substituted for leather or may behave like putty. The essential feature about soft matters, generally referred to as gels, is the existence of networks generated by chain couplings at widely separated loci in space; consequently, physical inquiry into the behavior of the statistical structures of polymeric networks must focus both on segmental properties corresponding to a few nanometer scale in space and on local properties induced by the dynamics of ten or less monomeric units. It is shown here, that in contrast to usual local nuclear magnetic resonance (NMR) approaches [1, 2] the proton magnetic relaxation is also a suitable technique for providing a nanometer scale characterisation of polymers observed above the glass transition temperature (Tg) [3]. This Chapter is intended to provide a simple guide for rapidly characterising soft polymeric matters; the characterisation relies on low resolution proton magnetic relaxation which leads to the determination of standard parameters unambiguously interpreted. Considering investigations into properties of soft polymers, the specificity of these physical systems, compared with ordinary liquids formed from small molecules, originates essentially from the linkage of chemical units whatever their nature; resulting chains comprise of thousands of bonds forming a backbone. Significant properties induced by the effect of linkage are, on the one hand, the curvilinear orientational correlation of units which determines a persistent length along one chain and gives rise to the stiffness effect. On the other hand, there is an elastic component which characterises the response to a stress and which originates from the presence of topological constraints: molecules having their atoms bonded in concatened sequences exhibit a physical exclusion of volume which precludes the passage of one chain directly through the backbone of another. The macromolecules of the liquid do not have time to adjust their positions by viscous movement over the interval during which a force is applied and the only possible motion is that of deformation analogous to the behaviour of a solid. There is, in turn, a resistance

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Spectroscopy of Rubbers and Rubbery Materials to shear forces; the higher the frequency of the stress the stronger the resistance. This effect adds to a natural component of viscosity due to the friction of chemical units, i.e., to the diffusion of molecular momentum. Both the atomic connections and the topological hindrance generate a frequency dependent effect of elasticity in macromolecular liquids; this effect is enhanced when polymeric networks are formed [4]. Both the elasticity and the viscosity components can be detected from the magnetic relaxation of nuclei attached to the polymer.

8.2 Polymeric Networks A wide variety of network structures are currently observed.

8.2.1 Molten High Polymers It is known that any high polymer is a non-Newtonian liquid which exhibits unusual flows: it climbs a rotating rod or it swells when it emerges from a tube [5]. Correspondingly, any molten high polymer conceals a temporary network which originates from the existence of an uncrossable backbone contour characterised by a well-defined mean length [6]. Furthermore, it is now well-established that the dynamic responses of molten polymers originate from the collective motions of segments along one chain; depending on the chain length, the time evolution of internal chain fluctuations exhibits characteristic exponents. The relevant space scale of description of chain dynamics is determined from the specific backbone contour length which governs the component of elasticity associated with the presence of a temporary network.

8.2.2 Crosslinked Chains Statistically defined structures may also arise from the formation of crosslinks in a melt; the resulting gels are described within a percolation framework which predicts the existence of definite meshes [7, 8]. Contact-lenses, jellies or even jellyfish are common examples of gels. Latex beads with specific functionalities attached, such as antigens, are used in biodiagnostics.

8.2.3 Semi-crystalline Polymers In semi-crystalline polymers, small ordered domains are connected to one another by chain segments; thus they are in coexistence with amorphous domains and their random

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NMR in Soft Polymeric Matter: Nanometer Scale Probe distribution in space can give rise to three-dimensional networks which originate complex textures [9] (Figure 8.1). The semi-crystallisation of polymers yields a soft matter which is currently substituted for leather.

8.2.4 Block Copolymers Similarly, a polymeric medium characterised by strong cohesion, is also obtained from di- or tri-block copolymers made by linking two or three chemically homogeneous sequences which are incompatible with one another; usually, the Tg of one of the two sequences is above room temperature while it is below for the other sequence [10]. There is a phase separation; glassy segments are connected to one another by amorphous segments and they play the role of ordered domains formed in semi-crystalline polymers.

8.2.5 Loaded Polymers The process of chain adsorption on mineral aggregates immersed in a molten polymer generates loops and tails on the surface of the filler; a chain structure thus appears. In

Figure 8.1 Topological constraints and defects along the chains hinder the complete crystallisation of polymers. Chain segments are ordered inside small domains; both the free enthalpy of bulk crystallisation and the surface energy are involved in the formation of domains which occur, consequently, at temperatures lower than the melting temperature

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Spectroscopy of Rubbers and Rubbery Materials addition to the adsorption effect, a process of connection of aggregates, bonded by the polymer may occur, leading to a percolation effect [11, 12]. Networks resulting from both the formation of sulfur crosslinks and the adsorption of polybutadiene chains on carbon black play a crucial role in the viscoelastic behaviour of tires.

8.2.6 Aggregated Polymers When considering structural aspects of polymeric systems, solutions wherein partial polymer association occurs, must also be taken into consideration. In concentrated or semi-dilute solutions, long polymer chains can form networks through the association of short segments randomly distributed along the chains; the physical association may arise from charge transfer or from hydrophobic interactions; networks may also result from the presence of chains which both enter in the formation of small aggregates and connect them to one another.

8.2.7 Network Distribution Function Any polymeric matter, observed above the Tg, can be pictured as an ensemble of chain segments issuing from coupling junctions, formed by entanglements, crosslinks, crystallites or contact points of chains adsorbed on the surface of mineral fillers. The formation of such junctions, also called nodes, gives rise to networks that are defined statistically in space; their structures consist of meshes determined by chain segments embedded in a viscous medium. The relevant linear space scale of description of the network structures is not the size of one monomeric unit; it is defined by about the mean distance between two consecutive nodes: ≈ 100 Å. The need for characterisations of network structures defined on a nanometer scale clearly arises from the description of physical properties of polymers observed above the Tg. Nodes are sensitive to macroscopic deformations of networks; they undergo displacements when a network is strained. The density of nodes is a crucial factor in determining the responses to macroscopic deformations, which may result from the effect of swelling induced by a good solvent or from a stress applied to the polymeric system. It is considered that the understanding of the behaviour of networks relies on the knowledge of the distribution of the ensemble of segments which connect nodes; let G(

< rij > ) denote the probability distribution function in which is the mean vector ζ ij

between two consecutive i and j nodes. The correlation length ζij depends on the number

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NMR in Soft Polymeric Matter: Nanometer Scale Probe of bonds nij in the segment that joins these two nodes; it reflects the statistical framework applied to the description of the polymer matter [7]. The statistics are specific to the state of the polymeric medium, which is observed. The existence of a distribution function cannot be detected directly but its moments are readily determined according to specific experimental procedures. It is shown, here, that the transverse magnetic relaxation curves of protons attached to strands can be given a simple analysis that reveals properties specific to the distribution function G(

< rij > ). ζ ij

8.3 Basis of the NMR Approach Any strong steady magnetic field applied to a spin-system induces naturally an axial symmetry for magnetic properties of nuclei; this axial symmetry leads to well distinct irreversible behaviours of the longitudinal and the transverse components of the macroscopic magnetisation, respectively. The longitudinal relaxation implies a quasiresonant exchange of energy between the spin-system and the molecular thermal bath: the frequency window of observation of random molecular motions is necessarily defined around the Larmor frequency of nuclear spins (≈ 109 rad.s-1). The transverse relaxation reflects mainly the loss of quantum phase coherence of spins and no displacements of nuclei are necessarily involved in such a process [13]. The axial symmetry is of particular interest for observing the broad spectrum of relaxation rates related to the hierarchy of fluctuations, which affect any polymer chain in a melt (from about 109 down to less than 1 s-1). For protons attached to polymer chains, the irreversible dynamics of the component parallel to the direction of the magnetic field is sensitive to properties generated by the local viscosity, which governs the random rotations of monomeric units, in the polymeric medium. With regard to the transverse magnetisation, the relaxation process cannot be analysed without considering the time interval allotted for the full random rotations of chemical units; it is close to the time interval (≈ 1 s or more) required for the full renewal of the chain configurations. However, this is too long a process for inducing any magnetic relaxation mechanism; consequently, the transverse component is sensitive to a part of the hierarchy of chain fluctuations, only. In other words, random motions of units are detected as non-isotropic rotations; the irreversible dynamics of the magnetisation is thus governed by the non-zero average of spin-spin interactions that results from the anisotropy of rotations of skeletal bonds, generated by topological constraints. A solidlike behaviour of the transverse magnetisation is expected to be associated to the property of elasticity, provided the time scale of the renewal of configurations is longer than the NMR scale of observation. This effect is considerably enhanced when permanent networks are observed.

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8.3.1 Chain Elongation The NMR approach relies on the sensitivity of this technique to the effect of reduction of entropy, which accompanies the formation of networks. Considering first a liquid made from small molecules, these undergo random rotations and translations, which manifest large amplitudes. Then, the existence of connections between atoms in a macromolecule restricts greatly the configurations these atoms may assume although angles of rotations about bonds of the molecular backbone permit a diversity of configurations; bond rotations are affected by internal potential energies of interactions between chemical groups. Thus, the linkage of chemical units forming macromolecules represents the first step in the entropy reduction. In any polymeric network observed above the T, the existence of coupling junctions, whatever their exact nature, eliminates translations and prevents the strands from completely reorientating; this in turn leads to the second step in the reduction of entropy. The natural tendency of a chain segment, comprised of nitrogen bonds, is to retract, at a given temperature; for such a chain segment, the equation of elasticity is derived from the contribution of the configurational entropy to the chain free energy: f = - 3kT

r Nb 2

where - f is the force exerted to the chain segment ends for counterbalancing the retractive force resulting from the small extension r (Figure 8.2).

Figure 8.2 Stretching vector of a chain segment. The schematic illustration shows that configurational fluctuations pervade a large volume

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NMR in Soft Polymeric Matter: Nanometer Scale Probe N is the number of skeletal bonds in one chain and k is the Boltzmann constant while b is the mean skeletal bond. The mean square end-to-end distance, Nb2, is also referred to as the square of the Gaussian correlation length between the chain ends, σ(N); it reflects the effect of linkage of chemical units. The previous relationship between the force f and the extension r is extended to any real chain submitted to a small elongation provided the correlation length, σ(N), includes the stiffness property of the polymer: σ(N)2 = λKNb2; λK is referred to as a persistence length. The related reduction of entropy is expressed as:

r2 ΔS = -3k . The longer the stretching vector, r, the higher the reduction of entropy. 2σ 2 Turning our attention to NMR, one chain segment is considered; its supposed fixed ends prevent the segment from completely reorientating: the freedom of internal bond rotations is impeded by the hindrance generated by the fixed ends; this constraint gives rise to fast but non-isotropic random rotations of skeletal bonds which induce, in turn, a non-zero average of the energy of magnetic interactions between protons attached to the segment. There are residual dipole-dipole interactions, hHR, expressed as:

hH R ≈ hH D

(3 cos2 θ r − 1)r 2 Λ N2 b2

the Hamiltonian, hHD, represents the strength of dipole-dipole interactions established in the absence of any chain fluctuations; the chain stiffness is accounted for by the parameter Λ and θr is the angle that the stretching vector, r, makes with the steady magnetic field direction [3]. This key equation ensures the transfer from local (a few Angströms) to semi-local (a few nanometers) NMR properties. The longer the stretching vector, the higher the strength of residual spin-spin interactions. Considering the residual interaction, hHR, the proton relaxation is thus governed by non-isotropic diffusional rotations of monomeric units which induce a solid-like behaviour of the transverse magnetisation; this is therefore sensitive in a specific way to the presence of chain junctions and consequently to the architecture of the polymeric network. In the frame rotating at the Larmor frequency of the spins (≈ 109 rad.s-1), the normalised transverse magnetisation associated to one segment is expressed as: mx(t) = Tr [exp( i HRt ) Mx exp(-i HRt) Mx )] /Tr [ Mx2)] Mx is the quantum operator associated with the transverse magnetisation. With regard to the proton magnetic relaxation, the probe determined by the end-to-end vector r, is substituted for any chemical unit attached to the chain segment. The observation is thus delocalised over the space scale defined by the distance r ≈ 5 nm. From the spectroscopy point of view, HD represents a dispersion of non-coherent broadening frequencies and

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Spectroscopy of Rubbers and Rubbery Materials the rate of fluctuations of this dispersion is that of internal bond rotations (≈ 109 rad.s-1): it is much higher than the width HD (≈ 105 rad.s-1); consequently, the observed residual interactions do not result from bond fluctuations which might be too slow, compared with the NMR scale (≈ 10-5 s) but they arise mainly, from the anisotropy of bond rotations.

8.3.2 NMR Evidence for Networks: Pseudo-solid Spin-echoes The NMR characterisation of polymeric systems requires first the search for the existence of networks. The observation of a time reversal effect, specific to residual spin-spin interactions, gives evidence for the presence of polymeric networks [3]. This property is reflected by so-called pseudo-solid spin-echoes formed by applying a suitable radiofrequency pulse sequence that results in a rotation of the spin operators (Figure 8.3). The partial recovery of the quantum phase coherence of nuclear dipoles originates from the non-commutative property of the Zeeman energy with the quantum operator which represents the residual interaction after rotating the spins. This rotation has no effect on the magnetisation dynamics when the residual interaction, hHR, is equal to zero. No

Figure 8.3 Normalised proton transverse relaxation curve recorded from end-linked calibrated chains. Eight pseudo-solid spin-echoes are recorded to illustrate the time reversal effect specific to the presence of any polymeric network Reproduced with permission from J.P. Cohen Addad and H. Montes, Macromolecules, 1997, 30, 12, 3678, Figure 6A. Copyright 1997, American Chemical Society

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NMR in Soft Polymeric Matter: Nanometer Scale Probe echoes are observed on ordinary liquids or on short polymer chains. The pulse sequence applied to the spin system for observing a pseudo-solid spin-echo is determined by: 90°/x - τ/2 - 180°/x - τ/2 - 90°/y - (t-τ)/2 - 180°/y - (t-τ/2).

8.4 Crosslinked Chains Permanent networks formed from crosslinked chains are of particular interest for illustrating the NMR approach. In this Section, it is shown how the proton transverse relaxation observed on networks is characterised in a standard way and provides values of relevant parameters.

8.4.1 End-linked Calibrated Chains Calibrated gels are synthesised by using polymer precursors also called telechelic chains because their ends bear chemical functions; they are characterised by a polydispersity index close to one. These chains are end-linked by reacting with a suitable chemical reagent. The gelation process is usually supposed to be near completion; correspondingly, the fraction of polymer extracted by washing samples is smaller than 2%. Since the length of segments is constant, the only relevant variable is the vector, r, between two consecutive nodes. The transverse magnetisation observed over a whole dry gel, is written as:

M xR (t) = ∫ Tr o [exp(iH R t )o M x exp( − iH R t)M x ]

1 3r 2 3 exp( − )d r (2πσ 2 )3 / 2 2σ 2

where the probability distribution function of the vector r is Gaussian. Taking the expression of HR into consideration, it is clearly seen that the reduced variable used to carry out the above integral is actually r/σ. Consequently, the effect of the non-zero average is accounted for by simply dividing HD by N, the number of bonds in any segment. The description of the effect of linkage on NMR amounts to considering the renormalised interaction defined by HD/N (≈ 103 rad.s-1). Correspondingly, the timescale of relaxation, currently observed, is about 1 ms for N = 100.

8.4.2 Characteristic NMR Rates Experimental transverse relaxation curves cannot be usually described from exponential time functions but integral treatments of these curves yield standard parameters equivalent

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Spectroscopy of Rubbers and Rubbery Materials to relaxation rates. Without entering into too many details, it may be worth noting that the integral treatments amount to calculating several moments of the probability distribution function of the end-to-end vector, r. It is more convenient to give the numerical estimate of moments of the distribution function than to determine the exact expression of the relaxation function M0m. The first and the third moments of the distribution function, G, are obtained from the two following integral treatments of the experimental relaxation curve:

φ1 = ∫ oM Tx (t)/ t dt 0

and ∞

φ3 = ∫ 0

dM Tx (t) / t dt dt

with MTx(0) = 1. It is shown that the quantity called φ1 is proportional to the first moment of the probability distribution function G(ρ) ∞

m1 = ∫ ρG( ρ )dρ 0

while the quantity called φ3 is proportional to the third moment of G(ρ) ∞

m 3 = ∫ ρ 3G( ρ )dρ 0

with ρ = /σ(N). More precisely,

φ1 ≈ Λ−1 / 2 Δ−G1 / 2 < σ ( N ) / b > m1 /m 2

(8.1)

φ3 ≈ Λ1 / 2 Δ1 / 2;G < b / σ ( N ) > m 3 /m 2

(8.2)

and

The quantum average Δ1/2;G characterises spin-spin interactions established in the absence of any chain fluctuations. The ratio defined by:

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NMR in Soft Polymeric Matter: Nanometer Scale Probe χ c = φ3 / φ1 is called NMR structural parameter; it is expressed according to Equations 8.1 and 8.2 as: χc ≈ Λ1/2Δ1/2;G m3/m2

(8.3) −1 / 2

and Δ*G is given by the ratio ( Δ G / Δ G ). The parameter χc is defined in a standard way and is like a relaxation rate; 1/χc is proportional to the correlation length of segments between two consecutive nodes. In contrast to the two- or three-dimensional NMR approaches, local information is sacrificed, here, in order to focus attention on the space averages that characterise the network structure considered as a whole. Integrals convert the time dimension of relaxation curves into simple numbers: it is a zero-dimension NMR approach. 1/ 2

8.4.3 Strand Length Dependence

8.4.3.1 Thermal Behaviour of Dry Gels Applying Equation 8.3 to calibrated dry gels, the correlation length between two consecutive nodes is expressed as σ(N)2 = λKNb2; 1/χc is thus proportional to the number of skeletal bonds, N, between two crosslinks. The thermal behaviour of the transverse relaxation exhibits three domains, usually observed on heating gels from the Tg(Mn) which is a function of the molecular weight of the polymer precursor, Mn. Chain configurations start taking place over the range Tg to Tg + 40 K but they fluctuate too slowly to be detected from NMR; the NMR response is thus insensitive to temperature variations. Then, the progressive onset of chain fluctuations as detected from NMR is observed over the range Tg + 40 K to Tg + 100 K. On raising the gel temperature, segmental fluctuations become faster and faster and more and more configurations are detected; there is a transient behaviour described by the empirical equation

1 = β [T − Tg (M n ) − T0 ][1 + γN ] χc

(8.4)

with β = 9 10 ms.K-1 and γ = 4.2 x 10-2; the quantity 1/χc is expressed in ms, in Equation 8.4. The last factor in Equation 8.4 is a linear function of the number of skeletal bonds, N, between two adjacent nodes; it reflects the effect of small elongation of network segments due to the presence of fixed ends. Finally, for temperatures higher than Tg + 100 K, configurational fluctuations are fully observed; they are only restricted by the small

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Spectroscopy of Rubbers and Rubbery Materials stretching of chain segments connecting any two nodes throughout the gel. Considering calibrated polypropylene-oxide gels, the relaxation rate has been shown to be simply proportional to 1/N: χc = γ p / N with γp = 23 ms-1; the molecular weight of segments was varied over the range 200 to 4000 [14].

8.4.3.2 Swelling Effect The dependence of NMR on the segmental length between crosslinks, is also conveniently detected from gels swollen by a good solvent. Chain segments obey the excluded volume statistics and the correlation length is written as σ(N)2 = N1.2 λKb2 [7]. The proton relaxation rate, χc, is then expressed as a function of the swelling ratio, Qm, as: χc ≈ Qm-1.5 Qm is the volume of the swollen gel divided by the volume of the dry gel: Qm ≈ N.0.8.

8.4.3.3 Linkage Effect Considering calibrated gels, the effect of linkage of chemical units on NMR is easily disclosed when chain ends are frozen by the formation of crosslinks. The linear molecular weight dependence of χc-1 may serve as a reference for calibrating other gels made from the same polymer species.

8.4.4 Randomly Crosslinked Chains Network structures are still determined by nodes and strands when long chains are crosslinked at random, but the segmental spacing between two consecutive crosslinks, along one chain, is not uniform in these systems which are currently described within the framework of bond percolation, considered within the mean field approximation. The percolation process is supposed to be developed on a Cayley tree [15, 16]. Polymer chains are considered as percolation units that will be linked to one another to form a gel. Chains bear chemical functions that can react with functions located on crosslinkers. The functionality of percolation units is determined by the mean number f of chemical functions per chain and the gelation (percolation) threshold is given by pc = (f-1)-1. The

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NMR in Soft Polymeric Matter: Nanometer Scale Probe probability that two chains have been bridged is called p. The variable of gelation (percolation) is then defined as the deviation from the gelation threshold ε = (p-pc)/pc. The formation of meshes accompanies the vulcanisation process; their size is a function of e and the resulting network structure is specific to the state of gelation associated to this variable. Thus, the weight fraction of synthesised gel, the modulus of elasticity, E and the swelling ratio, Qm are physical quantities which must depend on the variable of gelation. The transverse relaxation of protons attached to the chain segments is sensitive to the relevant space scale determined by the mesh size.

8.4.4.1 NMR-Vulcanisation Relationship More precisely, the magnetic relaxation depends on the variable of gelation, i.e., the density of crosslinks, and is closely related to the modulus of elasticity, E, on the one hand and to the swelling ratio, Qm, on the other hand. Long polybutadiene chains are currently randomly crosslinked, using sulfur; they can serve to illustrate the NMR approach to the characterisation of vulcanised polymers. It has been shown that the variable of gelation is expressed as: ε =

2γ sM W < ms >

γ s is the sulfur concentration (w/w),

<ms> is the mean weight of one sulfur bridge and Mw is the weight average polymer molecular weight. The probability that two polybutadiene chains have been bridged by sulfur is p=2γsMm/<ms> while the percolation threshold is defined as pc=Mm/Mw (Mm is the molar weight of one monomeric unit). The above equation relies on the assumptions that any monomeric unit can react with sulfur and that the sulfur is entirely involved in the formation of bridges between chains [17]. The standard relaxation rate, χc, is shown to be actually a function of ε when the weight average polymer molecular weight, Mw, is varied over the range 70 x 103 to 180 x 103 while the sulfur concentration, γs, is varied from 0.002 to 0.01 g/g. Similarly, the NMR-percolation relationship is also well illustrated from a slightly modified silicone which is a copolymer composed of dimethyl monomeric units (-O-Si(CH3)2) and of a small number of vinylmethyl units (-O-Si(CH3)(CH=CH2)), randomly distributed along the chains [18]. One vinyl group can react with a neighbouring methyl group, in the presence of a catalyst; the reaction occurs at 150 °C under pressure and yields a link between two chain segments. The kinetics of formation of crosslinks is observed by quenching polymer films in ice to stop the chemical reaction and to keep the polymer network in the gelation state reached at a given time t. The number of segments issuing from one node is 4. Let α(t) denote the fraction of vinyl groups that have reacted with methyl groups at a time t. The gelation variable ε is thus written as: ε(t) ≈ α(t)2 Cvi2 - 1,

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Spectroscopy of Rubbers and Rubbery Materials Cvi is the number of vinyl groups per chain. The direct observation of the decrease in amplitude of the Fourier Transform-Infrared spectroscopy (FT-IR) spectrum of vinyl groups leads to a square root dependence of α(t) on time. The threshold of gelation occurs at a time t0 such that ε(t0) is equal to zero. Again, the magnetic relaxation rate is a function of the variable, ε (Figure 8.4).

Figure 8.4 Typical variation of the standard NMR parameter, χc, as a function of the variable of percolation e; the different states of gelation were obtained by quenching the polymer in ice, at different times during the kinetics of synthesis of the network which occurs at 150 °C. The polymer is a slightly modified silicone chain which bears randomly distributed vinyl groups as comonomeric units (the mean number of vinyl groups along one chain is Cvi = 2 x 10-3). Links are created between vinyl and methyl groups; the three symbols correspond to different catalyst concentrations (redrawn from [18])

8.4.4.2 Elasticity-NMR Relationship One important feature of randomly crosslinked chains is that they can be highly strained and/or swollen; they recover their initial shape when the stress is interrupted and/or after deswelling. Neither the dependence of the modulus of elasticity, E, upon the variable ε nor that of the swelling ratio, Qm, have been exactly predicted, until now. Nevertheless, these two physical quantities as well as the rate χc are functions of the variable, ε.

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NMR in Soft Polymeric Matter: Nanometer Scale Probe Consequently, they are closely related to one another. The NMR-elasticity relationship is illustrated from the linear dependence of the relaxation rate, χc, on the modulus of elasticity of the gel, E (≈ 0.5 MPa): χc = χ0c+ qE with q ≈ 0.3 ms-1MPa-1, in Figure 8.5, corresponding to polydimethylsiloxane chains crosslinked at random and already described in Section 4.4.1; the relaxation rate, χ0c , measured in the absence of any strain, reflects the elasticity of the highly entangled molten polymer [19]. It is worth noting that NMR measurements are performed without any stretching of the polymeric gel. Segmental fluctuations involved in the linear mechanical response to a small strain are directly observed from the transverse magnetic relaxation, in the absence of any strain.

Figure 8.5 Linear dependence of the NMR standard parameter, cc, on the modulus of elasticity E measured on randomly crosslinked polydimethyl siloxane (PDMS) chains. NMR measurements were performed without any sample deformation (redrawn from [19]) Adapted with permission from J.P. Cohen Addad, B. Phan Thanh and H. Montes, Macromolecules, 1997, 30, 15, 4374, Figure 2. Copyright 1997, American Chemical Society

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8.4.4.3 Swelling Effect Similarly, considering the swelling effect of vulcanised poly(butadiene), induced by a good solvent, the relaxation rate has been shown to obey the simple equation: χc ≈ Qm-2 the dependence upon the swelling ratio, Qm, reveals the Gaussian statistics which applies to the description of the swollen system. It is contrasted with the Qm-1.5 dependence which reflects the excluded volume statistics specific to the description of end-linked calibrated chains swollen by a good solvent. Considering calibrated gels or randomly crosslinked chains, the effect of linkage of monomeric units on NMR is well detected in spite of the purely statistical definition of the segmental spacing between crosslinks in vulcanised chains.

8.4.4.4 Stretched Gels The anisotropy of random motions of monomeric units can be mechanically induced by stretching gels or vulcanised chains. The resulting non-zero average of the rank 2 tensor of space variables concerns the dipole-dipole interaction as well as the quadrupolar one. The deuterium NMR spectroscopy has been extensively studied; in particular, the quadrupolar splitting has been quantitatively related to the deformation ratio, λ, resulting from the uniaxial stretching of a network [20, 22]. The transverse proton magnetisation is also sensitive to the elongation of polymeric networks even though the interpretation of experimental results requires a careful analysis of the relaxation curves [23, 17].

8.4.5 Latex Suspensions The NMR approach applies to suspensions formed from latex beads. The micrometer particles are subject to Brownian motion from the thermal fluctuations in the solvent; however, the random rotation of beads is too long a process to induce any magnetic relaxation process. Consequently, particles can be considered as fixed and the transverse magnetic relaxation of protons attached to the latex reflects unambiguously network properties inside polymeric beads. This approach holds as long as the mean size of particles, D, is such that the rotational correlation time derived according to the Stokes equation (≈ 6πD3η0/kT, the solvent viscosity, η0 is set equal to 0.1 Ps-s and T = 300 K) is longer than the inverse of the residual proton-proton interactions: 10-3 s, i.e., D ≥10 nm.

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NMR in Soft Polymeric Matter: Nanometer Scale Probe

8.4.6 Kinetics of Gelation NMR is a suitable technique for characterising the kinetics of gelation of polymers, in situ. The sensitivity of the detection is considerably enhanced by observing amorphous segments rather than crosslinks. For example, the presence of only one crosslink is reflected by about 100 monomeric units located around it because they are submitted to additional topological constraints generated by the formation of one bridge between two segments [24].

8.5 Polymeric Crystallisation As one might anticipate, the transverse relaxation is greatly sensitive to the degree of crystallinity of a polymer. It has proved to be a suitable tool for discriminating the relaxation of protons embedded in solid domains from the relaxation of protons located in amorphous parts, in a semi-crystallised polymer. More precisely, in the case of semicrystallised polymers, ends of amorphous segments are embedded at the interface of the nuclei that form and which act like physical crosslinks. Fluctuations that occur along segments, with fixed ends, are conveniently characterised from NMR and the effect of confinement, exerted on amorphous segments by ordered domains can be examined. For a given state of semi-crystallisation of polymers the crystallinity can be correlated to the properties of proton relaxation in the amorphous state.

8.5.1 Crystallisation-NMR Relationship A great deal of effort has been devoted to characterising the crystallisation of pure polymers [25, 26]. Investigations have been extended to random copolymers, using x-ray scattering, differential scanning calorimetry (DSC) and polarised light microscopy [27, 28]. In this Section, attention is focused on the NMR characterisation of the state of semicrystallisation observed during the progressive annealing of propylene-ethylene copolymers, quenched at room temperature, from the melt. It has been shown that statistical copolymers with low ethylene contents crystallise from the melt in different forms. The development of each form depends on the cooling rate from the melt while the supermolecular structure is determined by the undercooling temperature of observation [29]. Ignoring the crystalline forms, the degree of crystallinity, Xc, can be detected, step by step, observing the relaxation of the transverse magnetisation of protons attached to the chains. The sharp decay associated with protons located in ordered domains is contrasted to the long decay corresponding to protons attached to amorphous segments. The ethylene content in random copolymers, hereafter called C2, varied over the range 0 to 4.7 % (w/w). Let T0m denote the melting temperature of pure polypropylene, the

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Spectroscopy of Rubbers and Rubbery Materials presence of co-monomeric units is expected to induce a depression of the melting temperature of the copolymer Tm(C2). The empirical temperature dependence of the degree of crystallinity, as detected from NMR, is found to obey a single curve; this curve is translated along the temperature axis in accordance with the melting point depression when the ethylene content varies: Xc (Tm(C2)-T). The degree of crystallinity is primarily a function of the undercooling, whatever the ethylene content. Considering any state of semi-crystallisation, a strong correlation between the degree of crystallinity and the relaxation rate of protons, attached to amorphous segments, is established. It is seen from Figure 8.6 that the NMR sensitivity to the crystallinity is considerably enhanced when detected from amorphous properties [30].

8.5.2 Kinetics of Crystallisation With regard to the crystallisation, polymers stand in sharp contrast to other materials such as metals that crystallise completely at the melting temperature, Tm; the topological

Figure 8.6 Variations of the superposition factor s with the degree of crystallinity Xc; s is the factor applied to the time scale to bring all relaxation curves into coincidence with a given curve chosen as a reference. ■ : pure polypropylene; ethylene content, ● : 4.7 % , ❍ : 3.8 % , ▲ : 3.7 % and Δ : 3.0 %, w/w (redrawn from reference [30]) Reproduced from L. Dujourdy, J.P. Bazile and J.P. Cohen Addad, Polymer International, 1999, 48, 561. Copyright Society of the Chemical Industry. Reproduced with permission. Permission is granted by John Wiley and Sons on behalf of the SCI

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NMR in Soft Polymeric Matter: Nanometer Scale Probe hindrance and the presence of chain microstructure defects prevent polymers from completely crystallising. Consequently, both the free enthalpy of bulk crystallisation per unit volume - Dhf, and the surface energy, γc - are involved in the formation of small ordered domains. This process occurs necessarily below the melting temperature and its kinetics (up to a few months) depend strongly on the undercooling temperature interval, (Tm-T). The curves of kinetics of isothermal crystallisation observed on most polymers can be analysed according to the equation proposed by Avrami [31]: Xc(t) =X ∞ [1- exp(-Ktn)] c Xc(t) is the crystallinity ratio measured at time t, while X ∞ is the maximum of the c crystallinity ratio measured at the temperature of observation, T; the exponent, n, depends on the mean shape of the ordered domains that are formed parallelepipeds or platelets. The analysis of the experimental curves of crystallisation kinetics leads to the determination of the kinetics constant, K which depends on the ratio γcTm/Δhf(Tm-T); the experimental value of γc is thus derived from K. This ratio is an estimate of the minimum size that nuclei which form spontaneously must have for starting to grow and giving rise to the polymer crystallisation. This analysis also applies to the crystallisation of statistical copolymers which takes place in solutions: the mean segmental spacing between two consecutive comonomeric units along one chain, required for the formation of ordered domains, must be longer than the length determined by the ratio γcTm/Δhf(Tm-T).

8.6 Entangled High Polymers The main feature about molten high polymers (molecular weights higher than about 104) concerns the broadness of the relaxation spectrum that characterises the viscoelastic response of these systems. This broad two-dispersion spectrum may spread over a range of relaxation times going from about 10-9 up to several seconds [4]. It is well illustrated from the modulus of relaxation observed after applying a sudden stress to the polymer; the resulting sudden deformation of the sample is then kept constant and the applied stress is released in order to avoid the flow of the polymer. For example, the release of the constraint σxy(t) is expressed as a function of the shear modulus of relaxation Gxy(t): Gxy(t): σxy(t) = Gxy(t)

where

∂γ xy ( t − > 0) dt ∂t

∂γ xy ( t − > 0) is the time derivative of the initial shear applied to the sample ∂t

during δt seconds. The time dependence of Gxy(t) is currently considered as reflecting the

309

Spectroscopy of Rubbers and Rubbery Materials actual dynamics of internal fluctuations of any chain, in the melt; it consists of three parts. For short times, t ≤ 10-5 s, Gxy(t) is a decreasing function, independent of the chain molecular weight and reflects the dispersion of relaxation times associated with segmental fluctuations. There is then a plateau, which corresponds to the break in the relaxation spectrum; it is characterised by a modulus of temporary elasticity G0N, about equal to 0.1 MPa. The plateau is involved in the storage of energy during any deformation, a crucial property for polymeric materials currently pictured as temporary networks; by analogy with the description of the property of elasticity of permanent gels, the equivalent mean number Ne of skeletal bonds in strands is expressed as: Ne = ρRT/G0NMb where ρ is the polymer density and Mb is the molar weight of one skeletal bond [4]. Finally, the second dispersion of relaxation times is associated with the full renewal of chain configurations which is achieved during the third part of the time dependence of Gxy(t); this relaxation, is characterised by a terminal rate TR-1 which depends strongly on the chain molecular weight, M: TR ≈ M3 [7]. The existence of a two-dispersion spectrum is an effect resulting from the linkage of monomeric units.

8.6.1 Temporary Networks It is known that rates of molecular fluctuations are easily detected from the proton magnetic relaxation only if they are higher than the spin-spin interactions (in rad.s-1) [13]. As a consequence of this general requirement, any molten high polymer, observed from NMR, behaves like a gel when the rate of dissociation of the temporary network is smaller than the residual dipole-dipole interaction of protons attached to the chains [32, 33]. The existence of a break in the relaxation spectrum implies that long-range chain fluctuations cannot be detected while segmental fluctuations are easily studied. Most observed transverse magnetic relaxation curves of protons exhibit, well shaped pseudosolid echoes which cannot be distinguished from echoes recorded from permanent gels; their observation gives a clear evidence for the existence of a temporary network as detected from NMR. In addition to this evidence, the maximum of the spin-lattice relaxation rate of protons, observed on high polymers, occurs at about 100 K above the Tg when the Larmor frequency is 109 rad.s-1; the maximum reveals, without any ambiguity, the presence of random motions of monomeric units characterised by a correlation time close to 10-9 s. The spin-lattice relaxation rate is usually independent of the chain molecular weight, except for short chains; in other words, long range effects are screened when observing local motions in strands determined by the mean number Ne of skeletal bonds. These motions are non-isotropic and give rise to the residual dipole-dipole interactions, which are observed in high polymers as well as in permanent gels. The pattern of transverse

310


Figure 8.7A: Semi-logarithmic plot of the proton transverse relaxation functions recorded from 1,2 polybutadiene, at 348 K. Vinyl contents are: χ1,2 = 0.82 (curve a), 0.66 (curve b), 0.58 (curve c), 0.40 (curve d) and 0.22 (curve e) 8.7B: Time scale shifts are applied to the relaxation functions to illustrate the property of superposition. Time shift factors are 1.05 (χ1,2 = 0.40), 1.16 (χ1,2 = 0.58), 1.30 (χ1,2 = 0.66), and 1.62 (χ1,2 = 0.82)

relaxation curves recorded from polybutadiene in the melt illustrates the existence of a temporary network as detected from NMR (Figure 8.7). In this system, both the Tg and the modulus of temporary elasticity G0N depend strongly on the concentration of monomeric units in the vinyl 1-2 conformation, along one chain.

311

Spectroscopy of Rubbers and Rubbery Materials The pattern can be obtained from the polymer temperature or concentration variations in addition to the change of G0N. The relaxation function may be too complicated a mathematical expression ever to be calculated, nonetheless, it obeys a property of invariance which allows the superposition of all normalised relaxation curves to one another by adjusting a suitable factor to the time scale of each curve. The time shift factor is found to obey the equation

s[T,Tg (φ ) − T0 ] ≈ [T − Tg (φ ) − T0 ][1 +

Γ ] G φα 0 N

(8.5)

where Γ is a constant; Tg(φ) is the Tg of the polymer solution and φ is the polymer concentration [34, 35]. The temperature T0 corresponds to the onset of fluctuations as detected from NMR; the exponent α is equal to 2.2. Equation 8.5 applies to polyisobutylene chains in concentrated solutions or to polyethylene-oxide, too. It is worth noting that the behaviour of the relaxation time scale, represented by Equation 8.5, extends to calibrated gels when the number N is substituted for the ratio Γ/G0Nφα ; this behaviour actually concerns protons attached to any chain segment embedded in a viscous medium.

8.6.2 Short Chain Dynamics The NMR approach is two-fold. Along with the characterisation of the equivalent network specific to any high polymer, this technique is a suitable tool for studying the dynamics of short chains (molecular weight < 10,000), in the melt. The principle of the description of the chain dynamics has been outlined 45 years ago; it is based on the assumption that information about the short-range behaviour at high frequencies can be sacrificed in order to reduce considerably the huge number of degrees of freedom, associated with any chain [36]. Any polymer is represented as a flexible chain of Gaussian submolecules and long range fluctuations are described as the collective motions of coupled submolecules; each Gaussian segment is submitted both to the effect of viscosity and to the property of elasticity which applies between its two ends. Any submolecule is characterised both by its time dependent end-to-end vector and by the related residual spin-spin interaction of attached nuclei which is in turn time dependent, too; coupled equations of all end-to-end vectors lead to the quantitative description of configurational fluctuations. The terminal chain relaxation time is proportional to the square of the chain molecular weight. Correspondingly, the slow time dependence of the residual interaction induces a transverse relaxation process specific to the long-range fluctuations of the chain. It has been shown that the molecular weight dependence of the proton relaxation rate is logarithmic [36].

312


8.6.3 Long Chain Dynamics Considering the hierarchy of long-range fluctuations that occur within long chains, the Rouse model applies also to the description of the collective motion of submolecules. During the time scale associated with the Rouse model, long range fluctuations are not sensitive to the back and forth displacement (reptation) of one chain supposed to move within a tube representing the lateral hindrance induced by surrounding chains [7]; the time scale associated with the reptation motion is much longer than the time scale determined by the Rouse model. This model was used as a convenient numerical tool for giving the quantitative interpretation of the proton transverse relaxation observed on polyethylene-oxide as a polymer molecular weight function [38, 39, 40]. It is now well established that the observed normalised relaxation function MTx(t) can be analysed using the equation:

M Tx (t) = (M xR (t) + M xEnd (t))Φ R (t) where ΦR(t) is the contribution to the relaxation induced by fast but non-isotropic rotations of monomeric units while MxEnd (t) is represented by an exponential function associated with the relaxation of so-called end-submolecules; finally, MxR(t) is the part of the relaxation function sensitive to the chain dynamics hindered by entanglements, with: u1 τs t /τs 1n(M (t))/t = − Δ τ s ( ) ∫ du1 ∫ du 2 Γ ( u 2 ) 0 t 0 R x

2

This second order expression accounts for experimental results with:

Γ( u ) = [

1 ∑ exp(− u ) / τp)]2 Ns p

according to this model, one chain is arbitrarily divided into Ns submolecules and the correlation time τp is defined as:

πρ 1 1 = sin 2 sm 2 Ns τ p τs

(8.6)

where τs is the correlation time associated with one submolecule comprised of n skeletal bonds:

313


τs =

C∞d 2ξ0 n 2 12 kT

(8.7)

Γ(|u|) is actually a function both of the single time constant τs and of the number of submolecules Ns in one chain; C∞ is the characteristic ratio of one polyethylene oxide (PEO) chain. The mean square length of one skeletal bond is d2 while ζ0 is the friction coefficient associated with one skeletal bond. The parameter Δ2 includes both the quantum and the spacial square averages of the spin-spin interactions resulting from local motions which occur within any submolecule; its numerical value cannot be exactly calculated but the order of magnitude of Δ, derived from the mean segmental spacing between entanglements, is 103 rad.s-1. The very good accuracy of the numerical description of the proton relaxation is illustrated in Figure 8.8 where the computed expression: [ln( MxT(t) )]/t is compared with the experimental curve recorded from polybutadiene observed at 318 K.

Figure 8.8 The computed expression (ln(MxT(t))/t is compared with the curve calculated from the relaxation curve recorded from molten polybutadiene (1.16 x 105 g/mol) at 318 K; = 1.05 rad.s-1, = 0.42 ms; Ns = 17

314

NMR in Soft Polymeric Matter: Nanometer Scale Probe The estimate of the numerical value of (Δ/Ns) amounts to dividing the dipole-dipole interaction, |HD| by the total number of skeletal bonds in one chain: N=nNs. This estimate (≈ 15 rad.s for a PEO chain molecular weight equal to 200 K) is independent of the partition of the chain into submolecules. The good agreement between the theory and experimental results provides the numerical value of the correlation time τs associated with one submolecule; then, knowing τs, the value of the friction coefficient ζ0 associated with one skeletal bond is obtained considering Equation 8.7. Furthermore, the NMR approach reveals without any ambiguity the existence of end-submolecules which play a crucial role in the reptation motion of one long chain in the melt ; the molecular weight of one end-submolecule is 11000 g/mol and 4850 g/mol for polyethylene-oxide and polybutadiene, respectively [39, 40]. The approach has been extended to PEO solutions.

8.7 Adsorption on Mineral Aggregates The accurate description of the mechanisms of reinforcement of elastomers is still a baffling problem which requires the characterisation of several phenomena involving mainly the polymer adsorption on the filler surface and the bond percolation resulting from the random polymeric connection of mineral aggregates. NMR applies to the determination of the law of chain adsorption on the filler surface by observing the relaxation of protons attached to monomeric units fixed on aggregates; knowing the weight of polymer adsorbed per gram of filler, this observation provides the mean number of contact points of one chain with the surface. Fumed silica is usually mixed with polydimethylsiloxane; it consists of aggregates with a fractal structure formed from sticked elementary beads (diameter ≈ 14 nm); the average length of aggregates is about 200 nm while their average diameter is about 70 nm and the specific surface area is 150 m2g-1. Silanol groups cover the silica surface (≈ 2 groups per nm2) and can form hydrogen bonds with oxygen atoms of siloxane chains. The mechanical mixing of the polymer with silica leads to a process of chain adsorption involving loops and tails which correspond to a configurational entropy reduction detected from NMR (Figure 8.9) ; both the initial amount of polymer in a mixture and the amount of adsorbed polymer on silica are conveniently detected from NMR [41]. The density of adsorbed chains is νc and the density profile of loops comprised of n units, S(n), has been predicted to vary as: S(n) = νc/n3/2 (N is the number of chemical units in one chain) [42]; the resulting mean number of contact points of one chain with the silica surface is N1/2 [11, 12]. The NMR sensitivity

315


Figure 8.9 Semi-logarithmic plot of the proton transverse relaxation curve recorded from a silica-polydimethylsiloxane mixture. The sharp decrease is associated with the proton magnetisation of adsorbed monomeric units while the long time dependence corresponds to loops and trains of linked chains

to the loop density profile is particularly enhanced by swelling the network using a good solvent; the corresponding transverse magnetic relaxation rate has been shown to vary as N-1/2. Such a specific dependence results from a purely structural property of adsorption and no chain diffusion process is involved in the proton relaxation process. A network is obtained when the relevant percolation variable, N1/2φSi, determined from the silica density, φSi, and N, has suitable values. The fraction of bound polymer, fB obeys the equation: fB = βN1/2φSi (1- βN1/2φSi /4 ) where β accounts for the specific surface area of silica AT ; β = AT Mm1/2/( σeA0 ) σe is the surface area associated with one silanol group, Mm is the molar weight associated with one skeletal bond and AO is the Avogadro number [11, 43]. Finally, NMR is a suitable tool for characterising the dynamics of short chains moving through the network formed by aggregates connected to one another by polymer chains [44].

316

NMR in Soft Polymeric Matter: Nanometer Scale Probe Similar approaches apply to carbon black - polybutadiene mixtures [45, 46]. A thorough study of ethylene-propylene diene terpolymer (EPDM) - carbon black has been developed using NMR; a clear evidence for adsorbed EPDM is given in addition to the presence of polymeric strands connecting aggregates to one another [47].

8.8 Conclusion Properties of soft polymers, neither solids nor true liquids, have been investigated, for many years, from their mechanical behaviour. The description of their texture has been mainly referred to their macroscopic responses to external stresses or strains whether these were static or time dependent. For several years, light scattering, fluorescence correlation, and neutron scattering experiments have been used to get a stronger correlation between the textures of networks and their mechanical behaviour. In this Chapter, attention was focused essentially on the simple and non-expensive low resolution NMR approach to the characterisation of polymeric network structures whatever their physical origin. Investigations are based on the observation of the transverse relaxation of protons attached to chain segments. In addition to the conventional effect of viscosity well observed from NMR in any liquid, the transverse relaxation is particularly sensitive to reductions of configurational entropy induced by topological constraints resulting from coupling junctions of chains. The effect of linkage of chemical units gives rise to a specific effect that underlies the NMR approach; it provides us with observations delocalised over more than 5 nm. Examples, chosen in this survey, highlight this characteristic property. Nowadays, the quantitative interpretation of properties of the proton transverse relaxation may lead to a sharp characterisation of complex rubbery materials, provided calibrated samples associated to each polymer species have been previously studied for determining standard NMR parameters. Here, the NMR technique is a suitable tool for observing the mesh size of any network structure including adsorption on mineral aggregates and the lifetime of the networks when molten polymers are considered. The deuterium and 13C NMR spectroscopies provide useful additional insights into the mechanisms of network deformations which are mechanically induced.

References 1.

R.R. Ernst, G. Bodenhausen and A. Wokaun, Principles of NMR in One and Two Dimensions, Clarendon Press, Oxford, UK, 1987.

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K. Schmidt-Rohr and H.W. Spiess, Multidimensional Solid-state NMR and Polymers, Academic Press, New York, USA, 1994.

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J.P. Cohen Addad, NMR and Fractal Properties of Polymeric Liquids and Gels, Progress in NMR Spectroscopy, Eds., J.W. Emsley, J. Feeney and L.H. Sutcliffe, Pergamon Press, Oxford, England, 1993.

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J.D. Ferry, Viscoelastic Properties of Polymers, 3rd Edition, John Wiley, New York, USA, 1980.

5.

R.B. Bird and C.F. Curtis, Physics Today, 1984, 37, 26.

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W.W. Graessley, Advances in Polymer Science, 1974, 16, 1.

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P.G. De Gennes, Scaling Concepts in Polymer Physics, Cornell University Press, Ithaca, NY, USA, 1979.

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D. Stauffer and A. Aharony, Introduction to Percolation Theory, 3rd Edition, Taylor & Francis, UK, 1994.

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L. Mandelkern, Crystallisation of Polymers, McGraw-Hill, New York, NY, USA, 1964.

10. F.S. Bates and G.H. Fredrickson, Physics Today, 1999, 52, 32. 11. J.P. Cohen Addad, Polymer, 1992, 33, 2762. 12. J.P. Cohen Addad, Polymer, 1989, 30, 1820. 13. A. Abragam, Principles of Nuclear Magnetism, Oxford University Press, Oxford, UK, 1961. 14. J.P. Cohen Addad, L. Pellicioli and J.J.H. Nusselder, Polymer Gels and Networks, 1997, 5, 201. 15. W.H. Stockmayer, Journal of Chemical Physics, 1943, 11, 45. 16. P.J. Flory, Principles of Polymer Chemistry, Cornell University Press, Ithaca, NY, USA, 1969. 17. J.P. Cohen Addad and E. Soyez, Macromolecules, 1992, 25, 6855. 18. J.P. Cohen Addad and H. Montes, Macromolecules, 1997, 30, 3678. 19. J.P. Cohen Addad, B. Phan Thanh and H. Montes, Macromolecules, 1997, 30, 4374.

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NMR in Soft Polymeric Matter: Nanometer Scale Probe 20. B. Deloche and E.T. Samulski, Macromolecules, 1981, 14, 575. 21. P. Sotta and B. Deloche, Macromolecules, 1990, 23, 1999. 22. A.I. Nakatani, M.D. Poliks and E.T. Samulski, Macromolecules, 1990, 23, 2686. 23. P.T. Callaghan and E.T. Samulski, Macromolecules, 1997, 30, 113. 24. J.P. Cohen Addad, E. Soyez and A. Viallat, Macromolecules, 1992, 25, 1259. 25. B. Wunderlich, Macromolecular Physics, Volume 3, Academic Press, New York, NY, USA, 1980. 26. B. Lotz, J.C. Wittman and A.J. Lovinger, Polymer, 1996, 37, 4979. 27. S. Laihonen, U.W. Gedde, P.E. Werner and J. Martinez-Salazar, Polymer, 1997, 38, 361. 28. H.J. Zimmermann, Journal of Macromolecular Science and Physics, 1993, B32, 141. 29. S. Laihonen, U.W. Gedde, P.E. Werner, M. Westdahl, P. Jääskeläinen and J. Martinez-Salazar, Polymer, 1997, 38, 361. 30. L. Dujourdy, J.P. Bazile and J.P. Cohen Addad, Polymer International, 1999, 48, 561. 31. M. Avrami, Journal of Chemical Physics, 1941, 9, 177. 32. J.P. Cohen Addad, A. Guillermo and C. Lartigue, Physics Review Letters, 1995, 74, 3820. 33. P. Sotta, C. Fülber, D.E. Demco, B. Blümich and H.W. Spiess, Macromolecules, 1996, 29, 6222. 34. A. Labouriau and J.P. Cohen Addad, Journal of Chemical Physics, 1991, 94, 3242. 35. A. Labouriau and J.P. Cohen Addad, Journal of Chemical Physics, 1991, 94, 3242. 36. P. Rouse, Journal of Chemical Physics, 1953, 21, 1272. 37. M.G. Brereton and M.E. Ries, Macromolecules, 1996, 29, 2644.

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Spectroscopy of Rubbers and Rubbery Materials 38. J.P. Cohen Addad and A. Guillermo, Journal of Chemical Physics, 1999, 111, 7131. 39. A. Guillermo, J.P. Cohen Addad and D. Bytchenk, Journal of Chemical Physics, 2000, 113, 5098. 40. J.P. Cohen Addad and A. Guillermo, Physics Review Letters, 2000, 85, 3432. 41. J.P. Cohen Addad and L Dujourdy, Polymer Bulletin, 1998, 41, 253. 42. M. Aubouy, Physics Reviews B, 1997, 56, 3370. 43. J.P. Cohen Addad and N. Morel, Journal Physique, 1996, 6, 267. 44. J.P. Cohen Addad and O. Girard, Macromolecules, 1992, 25, 593. 45. J.P. Cohen Addad and P. Frebourg, Polymer, 1996, 37, 4235. 46. V.J. McBrierty and K.J. Packer, Nuclear Magnetic Resonance in Solid Polymers, Cambridge University Press, Cambridge, UK, 1993. 47. V.M. Litvinov and P.A.M. Steeman, Macromolecules, 1999, 32, 8476.

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Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR

9

Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR Dallas D. Parker and Jack L. Koenig

9.1 Introduction It is in the vulcanisation process of elastomeric materials that the elastic properties are generated. The formation of crosslinks in rubbers prevents permanent deformation under load and ensures elastic recovery on removal of the load. Because the type of crosslinking process (sulfur, peroxide, radiation, etc.) and the number of crosslinks can radically change the physical properties of rubber, an understanding of the mechanism and structure of vulcanisation is extremely important. However, from a molecular characterisation viewpoint, the introduction of crosslinking limits the ability to characterise those systems because of the insolubility of the resulting network. In addition, the network structures formed during vulcanisation are complex and diverse in nature. Additionally, the extremely low concentration of the chemically modified structures induced by the vulcanisation approach the detection limits of many traditional analytical techniques. Many approaches including physical tests and chemical analysis have been attempted in an effort to relate the chemical microstructure to the physical properties of both raw and cured elastomers. Nuclear magnetic resonance (NMR) spectroscopy has become one of the most important analytical techniques used in the characterisation of materials. In the field of macromolecules, its use extends from monomer characterisation, through polymerisation kinetics and mechanisms, to direct observation of the chemical structures of polymeric materials [1]. Using NMR, information obtainable on polymer structures includes main chain microstructures (conformation, geometric isomerisation, spatial distances, etc.) comonomer composition and sequence, end and side group analysis, branching and crosslinking, abnormal structures (cyclic and isomerised structures), bonding, region enchainment, and tacticity. One advantage of NMR spectroscopy is that it can be used on almost all polymeric materials through either solid-state or solution NMR methods. However, the techniques and resolution of the two methods are radically different. For example, the proton NMR spectrum of the water is sharp and narrow with a bandwidth of 1 Hz, while the proton NMR spectrum of ice is extremely broad with a bandwidth of 20 KHz. The differences

321

Spectroscopy of Rubbers and Rubbery Materials in the NMR spectra of solids and liquids are due to a motional averaging of interactions [2, 3]. In liquids and solutions, local interacting fields are averaged to zero by the rapid isotropic motions of the nuclei (termed incoherent averaging) resulting in narrow linewidths. Anisotropic interactions, such as dipolar and quadrupolar interactions and chemical shift anisotropy, are averaged to zero by the molecular motions to effectively remove them from broadening the resonances in the observed spectra. In solution NMR, because of the lower frequency motions of polymers, incomplete averaging of anisotropic effects is observed to a certain extent resulting in some peak broadening. In solids, however, this effect is highly magnified. Because there is not sufficient motion to average the anisotropic interactions extremely broad lines are present often encompassing much of the entire spectrum. Because the incoherent averaging (molecular motion) does not narrow the NMR lines, coherent-averaging techniques such as dipolar decoupling (DD) and magic-angle spinning (MAS) must be used to produce narrow line widths [4]. In solids, the major contribution to 13C NMR line broadening is due to heteronuclear dipole coupling with attached protons. By using broad band irradiation in the proton field, the effect of C-H dipole decoupling is drastically reduced. In addition, rapid sample spinning at the ‘magic angle’ of 54°44´, causes the term (3cos2 θ –1) to become zero in the dipolar Hamiltonian, thus further decreasing heteronuclear interaction. In 1H solid state NMR, CRAMPS (Combined Rotation And Multiple-Pulse Spectroscopy) techniques combine MAS and multiple pulses (WAHUHA, MREV-8, BR-24, etc.) to reduce homonuclear dipolar interactions. The low sensitivity of the 13C nucleus is improved by transferring polarisation from the magnetisation-rich protons to the 13C nuclei by using cross-polarisation. By combining the techniques of high-power decoupling (DD), MAS, and cross-polarisation (CP) experiments into one grand experiment [5], narrow lines and enhanced sensitivity can be obtained for polymers in the solid state. As a consequence, high-resolution NMR spectroscopy has become an important tool in the structural investigation of polymers in the solid state [6, 7, 8, 9]. Recently, solid-state 13C NMR has been widely applied for the characterisation of vulcanised rubber systems [10] with several reviews published covering the solid-state NMR analysis of crosslinked elastomers [11, 12].

9.2 Sulfur Vulcanisation Mechanism Although sulfur vulcanisation was discovered over one hundred and fifty years ago, the exact mechanism of vulcanisation is still being examined. This arises not only from the complexity of the reactions and products formed but also to the fact that the mechanism of accelerated sulfur vulcanisation changes is dependent on the class of accelerators/ activators used. Typically, benzothiazole or sulfenamide are used as accelerators, zinc

322

Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR oxide as the activator, and a fatty acid such as stearic acid used as the coactivator. A generally accepted sequence of reactions is as follows [13, 14]: • An interaction of the curatives occurs to form the active sulfurating agent, Ac-Sx- Ac by a reaction of accelerator (Ac) and activator with sulfur: Ac + S8 → Ac - Sx – Ac [active sulfurating agent] • The polymer chains interact with the sulfurating agent to form polysulfidic pendent groups terminated by accelerator groups: Ac - Sx - Ac + RH → R - Sx - Ac + AcH [pendent sulfurating agent] where RH is the rubber chain. • Polysulfidic crosslinks are formed; R - Sx - Ac + RH → R - Sx – R + AcH [crosslinks] • Network maturing and competing side reactions also occur which do not lead to effective crosslinks. Thermal decomposition leads to the following reactions: R -Sx- Ac → cyclic sulfides + dienes + ZnS [degradation] R -Sx- Ac → R - S - Ac + Sx-1 [desulfuration] R - Sy - R → R - S - R + Sy-1 [monosulfidic crosslinks] R - Sx+y - R + Ac - Sz - Ac → R - Sx - R + Ac - Sy+z - Ac [sulfur exchange] During the curing and network maturing periods, there are at least three competing reactions: crosslinking, desulfuration, and degradation reactions. Desulfuration is a reduction in sulfur rank of the crosslinked chain while decomposition or degradation refers to a complete loss of the crosslink with formation of conjugated dienes, cyclic structures, thiols, etc. The network structures formed depend not only on temperature but accelerator types and concentration. The ratio of poly-, di-, and monsulfidic crosslinks strongly depends on the ratio of sulfur to accelerator in the formulation. Thus, the reactions of diene rubber with sulfur produce a variety of sulfurised structures. A schematic representation of different types of sulfurised structures in natural rubber (NR) is shown in Figure 9.1. The type of sulfurisation has been designated as A1, A2, B1, B2 and C1-types depending on the positions of sulfide attachment as seen in Figure 9.1 [15]. The c and t indicate cis and trans isomers of the structures. The main chain, saturated

323


Figure 9.1 Designation of nomenclature for structures occurring upon vulcanisation of NR and cis-polyisoprene

and unsaturated structures are also indicated. Crosslink structures may also be classified as efficient (intermolecular) or inefficient (intramolecular, cyclic, pendent) in nature.

9.3 The NMR Methods for Assigning Resonances to Chemical Structure One of the earliest studies of vulcanisate structure by solid-state NMR was the sulfur vulcanisation of NR with samples compounded under different formulations of accelerator and sulfur [15]. The 13C NMR spectra of NR samples cured with 10% sulfur at different cure times acquired with MAS and scalar decoupling is shown in Figure 9.2. Line broadening and decreased resolution are observed with longer cure times. A dominant new resonance was observed at 58 ppm that is assigned to the polysulfide or vicinal crosslink structure. The spectral resolution of the solid-state spectra was found to improve on application of high-power dipolar decoupling along with MAS [16]. With the increased resolution, new resonances were observed in this experiment upon vulcanisation at 14.0, 16.1, 18.0, 30.0, 40.1, 44.6, 50.2 and 57.5 ppm. The peaks at 40.1 and 16.1 ppm are generated from the C-1 and C-5 carbons of the trans-isoprene units. The increase in these two resonances with cure suggests the occurrence of cis-to-trans isomerisation during the curing process. Intensities of the peaks at 14.0, 44.6 and 57.5 ppm also increased with cure time. To aid in the NMR peak assignments of those 13C NMR spectra, the NR model compounds, 2-methyl-2-pentene, 2-methyl-1-pentene, and 4-methyl-2-pentene, were

324


Figure 9.2 Stack plots of 10% sulphur-cured NR for different times of cure. B: butadiene-like species, P; polysulfidic crosslinks Carbon numbers in 1,4-isoprene unit in NR have been designated as follows;-C1-C2(-C5)=C3-C4- or -Cγ-Cα(-Ce)=Cβ-Cδ-

sulfur-vulcanised and studied by solution 13C NMR spectroscopy [16]. In addition to the C NMR observations of the model compounds, the chemical shifts of the carbons for these structures were calculated. The aliphatic additivity constants for the sulfur groups were obtained from ‘The Sadtler Guide to Carbon-13 NMR Spectra’ [17] and added to the carbons of polyisoprene and ethylene-propylene copolymer. Figure 9.3 shows the calculated 13C chemical shifts of the model structures. These data indicate that the monosulfidic and disulfidic crosslinks should be distinguishable in the 13C NMR spectra of sulfur-vulcanised NR. However, it may not be possible to distinguish disulfidic crosslinks from crosslinks of higher sulfur rank (polysulfide). 13

325


Figure 9.3 The calculated 13C chemical shifts of the model vulcanised structures

The chemical shifts due to the monosulfidic crosslinks are influenced not only by the position on the monomeric unit to which it belongs, but also by the position of the carbon atom of the monomeric unit on the other side of the bridge. The shielding parameters of monosulfide substitution on the individual carbons of the isoprene unit have also been determined. It is shown that resolvable polysulfidic crosslink resonances exist in all positions of the backbone carbons while monosulfidic crosslinks appear only between C-1 and C-4 carbons with detectable intensity [18]. The complementary application of the NMR inversion recovery measurements (180o-τ90o-tR) and the computer fitting of the overlapping regions was found to be a useful method for structural analysis of vulcanised NR in the solid state [19]. Since the line widths in 13C NMR spectra of solids are relatively broad compared with the differences between chemical shifts, some weak signals are completely obscured in the resulting spectra. If the resonances have sufficiently different relaxation times, such as with methyl and methylene carbons, it is possible to detect neighbouring overlapped signals by using the inversion recovery delay τ value at which the interfering strong resonance has null intensity.

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9.4 Unaccelerated Sulfur-vulcanisation of NR There is wide variety of vulcanisation agents and methods available for crosslinking rubber materials including peroxide, radiation, urethane, amine-boranes, and sulfur compounds [20]. Because of its superior mechanical and elastic properties, ease in use, and low cost, sulfur vulcanisation is the most widely used. Although vulcanisation with sulfur alone is not practical compared to the accelerated sulfur vulcanisation in terms of the slower cure rate and inferior physical properties of the end products, many fundamental aspects can be learned from such a simply formulated vulcanisation system. The use of sulfur alone to cure NR is typically inefficient, i.e., requiring 45~55 sulfur atoms per crosslink [21], and tends to produce a large portion of intramolecular (cyclic) crosslinks. However, such ineffective crosslink structures are of interest in the understanding of complex nature of vulcanisation reactions. The solid-state 13C NMR investigation of unaccelerated, 10% sulfur vulcanisation at different cure times was performed [16]. Using analysis at 80 °C and swelling of the vulcanisate samples to increase chain mobility, major peaks were observed at 57.5, 56.5, 44.6, 43.0, 40.1, 38.6, 36.1, 30.0, 18.0, 16.1, and 14.0 ppm. The peaks at 57.5, 44.6, 40.1, 16.1, and 14 ppm all increase in intensity at approximately the same rate with increasing cure time while the peak at 50.2 ppm appears at early cure times and disappears after 30 minutes of cure. The remaining peaks show greater increase with increasing cure time. All NMR peaks show increased intensities with increasing sulfur contents of 10%, 20%, and 30% by weight. The peak at 50 ppm was assigned to the carbons involved in the polysulfidic crosslinks of the A1-type structure. The resonance peak, which appears at roughly 45 ppm in the spectra, was assigned to the polysulfidic crosslinks of B1-type structure although the monosulfidic A1-type crosslinks may also overlap with this resonance peak. The peak at 58 ppm is assigned to the crosslink point found in B1-type polysulfide structure. The splitting of this peak observed in the high temperature measurement was postulated to arise from differences in carbon configurations or conformations in the same type of structure. The small peaks found at 30.0, 36.1 and 38.6 and at the shoulder of 43 ppm, which were resolved in the high temperature measurement, were assigned to cyclic sulfides.

9.5 Accelerated Sulfur-vulcanisation of NR and IR Accelerated sulfur formulations are the most common vulcanisation systems used in commercial and industrial applications. Therefore, research on both the fundamental and applied aspects of accelerated sulfur vulcanisation is ongoing. Several reviews of the chemistry and/or physics of accelerated sulfur-vulcanisation of elastomers have been published [13, 14, 22, 23].

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Spectroscopy of Rubbers and Rubbery Materials The structure of NR vulcanised with sulfur contents of 0, 2, and 5 pph and the accelerators of tetramethyl thiuram disulfide (TMTD), N-cyclohexyl-2 benzothiazole sulfenamide (CBS), and N-oxydiethylene-2-benzothiazole sulfenamide (MOR) was studied by solid-state 13C NMR [23]. MOR is categorised as a benzothiazolesulfenamide and is more popularly called 2-(4-morpholinothio)benzothiazole (MBS). On examination of the spectrum generated by the sulfur donor TMTD and 0 pph added sulfur, no crosslink peaks were observed implying that the concentration of monosulfidic crosslinks was lower than the NMR detection limits. A significant amount of cis-trans isomerisation was noted by the increase in the trans isomer peaks at 16.1 and 40.1 ppm. On comparison of the TMTD vulcanised spectrum with unaccelerated sulfur vulcanisation from a previous study [16], many similarities were noted especially the formation of prominent peaks at 44 and 58 ppm. However, the accelerator generated spectra shows fewer peaks implying a simpler network structure with fewer side reactions. With accelerated systems, new resonances were observed upon vulcanisation at 15, 31, 34, 45, 58 and 130 ppm and at the shoulder of 30 and 41 ppm along with the trans isomers appearing at 17 and 41 ppm. Model compounds based on 2-methyl-2-pentene were studied to supplement the 13C chemical shift assignments of the products from accelerated sulfur vulcanisation of NR. It is observed in the model compound data that it may not be possible to distinguish between a 13C NMR resonance which is due to disulfidic crosslinks and a peak due to pendent accelerator groups, while a large chemical shift difference (~3 ppm) is observed for the monosulfidic bonds. The MBS-accelerated sample shows similar new resonances as seen in the TMTD accelerated systems. In this comparison however, the quantitative aspects of the data might be obscured due to the differences in the state of cure among the different formulations. The 13C chemical shifts were assigned in more detail for monosulfidic and polysulfidic crosslinks occurring in the accelerated sulfur vulcanisation of NR [18]. The NR was cured with a pure thiuram formulation (TMTD alone) in order to predominantly prepare monosulfidic bridges in the network. The distortionless enhancement by polarisation transfer (DEPT) experiments, in which the carbons with different level of protonation can be distinguished [22-24], were performed for the NR cured with extended levels of sulfur. Based on the DEPT results and previously reported model compound results [20], the chemical shifts of the resonances occurring in the spectra were assigned. In the DEPT experiments, both peaks around 50 and 58 ppm are divided into three components and the levels of the protonation for these six individual resonances are evaluated. The peaks at 57.4, 58.0 and 58.6 ppm are assigned to the polysulfidic crosslinks in the A1, B1 and B2-type structures, respectively. The peaks at 37.2 and 50.7 ppm are due to A1-type polysulfidic crosslinks. There was no apparent structural match for the quaternary peak at 50.2 ppm.

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Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR Another approach in the 13C NMR technique has been utilised to assign the peaks for the sulfurised structures occurring in the accelerated sulfur-vulcanisation of NR [25]. The NR formulations with a variation of accelerator (CBS) to sulfur ratio were cured at 150 °C to t90 (90% of the rheometer torque increase) plus 10 minutes. The Attached Proton Test (APT) experiment [13, 14] was performed for those vulcanisates, in which the carbon with even number of hydrogen (methylene or quaternary carbon) and the carbon with odd number of hydrogens (methine or methyl carbon) can be distinguished. In addition, the 13C spin-lattice relaxation times (T1) for the individual carbons in both backbone and sulfurised structures were measured. The carbons with odd number of hydrogens can be easily distinguished between methine and methyl carbons by observing the region where the resonance is occurring. The carbons that have even numbers of hydrogens with long T1 (1s~) are assignable to the quaternary carbons while those with short T1s are due to the methylene carbons. The resonances at 57.5 and 50.4 ppm are assigned to the polysulfide and monosulfide of B2-type structures, respectively. The peaks at 50.8 and 55.6 ppm are due to the polysulfidic bonds of A1 and B1-type structures, respectively, however the monosulfidic counterparts for these structures, which fit the APT results, are missing. None of the structures can explain the splitting of the peaks near 64 ppm. The T1 value of 1.1 seconds of these coupled peaks suggests that the resonances are due to oxidation products. The trans structures increase with increasing number of sulfide attachments to the polymer backbone. The peaks appearing at 12.9 and 14.3 ppm are possibly due to the end groups of hydrocarbon fragments formed during the mixing and crosslinking steps of sample preparation. The crosslink densities of the rubber networks were determined by the quantitative treatment of the 13C NMR spectra [26]. In this study, the peaks at 37.3, 44.7, 48.2, 49.6, 50.6, 52.5, 54.8 and 57.5 ppm are involved in crosslinks according to the peak assignments. One-half the sum of these areas should give the crosslink density of the sample as determined by 13C MAS NMR. The results were compared to the network chain densities obtained by n-heptane swelling methods to examine the efficiency of the intermolecular crosslinking reactions for the individual formulation. For the efficient cure system (EV); higher accelerator concentration relative to sulfur), nearly 100% of the sulfurisation is monosulfide and is involved in crosslinks. However for conventional cure systems or systems containing predominantly polysulfidic bonds, about 50% of the sulfide units are not involved in crosslinking. Solid-state 13C MAS NMR has been applied for quantitative determination of crosslink density in accelerated sulfur-vulcanised NR [27]. The concepts used to calculate the crosslink density by 13C NMR are the same as the one mentioned above, but different resonances were used for the quantitative treatment based on the different assignments

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Spectroscopy of Rubbers and Rubbery Materials of the 13C resonances [28]. The crosslink density of the network is evaluated by the following formulas: μC =

I(SX ) ρ ⋅ Io Mo

I(SX ) =

1 [I(58.6) + I(58.0) + I(57.4) + I(44.4) + I(44.5) + I(40.7)] 2

where Io is the reference intensity of the monomer unit, ρ is the density of the vulcanisates and Mo is the molar mass of the monomer unit. The peaks at 40.7 and 44.5 ppm are due to monosulfides and the others are assigned to polysulfides. If the overlapping peak around 45 ppm can be precisely decomposed to polysulfidic (44.5 ppm) and monosulfidic (44.4 ppm) contributions, the densities of poly- and mono-sulfidic crosslinks are individually determined by the 13C NMR. The 13C NMR crosslink density results were compared with the crosslink density obtained by the mechanical measurements. In the determination of the crosslink density by mechanical methods, the contributions of the topological constraints on the results were neglected and the density was expressed as G/2RT. The 13C and mechanical-crosslink densities were obtained for both sulfur and dicumyl peroxide (DCP)-cured samples to ensure the effect of wasted crosslinks (pendent or intramolecular type sulfurisations), which are expected in the typical sulfur-vulcanisation of NR. In the major range of crosslink densities, the crosslink densities for those two systems are described by the same linear function with a slope of 1.0. Based on these observations, it is shown that the crosslink density of the sulfur-vulcanised NR as determined by 13C is identical with the true crosslink density, and the influence of the wasted or ineffective crosslinks (pendent and cyclic crosslinks) and chain ends is negligible. However, this conclusion seems to be only valid if the effect of topological constraints or entrapped entanglements on the mechanical modulus is negligible which is rarely the case in real systems. The chemical microstructures of cis-polyisoprene (IIR) vulcanised with sulfur and N-tbutyl-2-benzothiazole sulfenamide (TBBS) accelerator were studied as a function of extent of cure and accelerator to sulfur ratio in the formulations by solid-state 13C NMR spectroscopy at 75.5 MHz [29]. Conventional (TBBS/Sulfur=0.75/2.38), semi-efficient (SEV=1.50/1.50) and efficient (EV=3.00/1.08) vulcanisation formulations were prepared, which were cured to different cure states according to the magnitude of increase in rheometer torque. The order and types of the sulfurisation products formed are constant in all the formulation systems with different accelerator to sulfur ratios. However, the amount of sulfurisation has been found to vary directly with the concentration of elemental sulfur.

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Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR Similar vulcanisation chemistry is observed with the N-t-butyl-2-benzothiazole sulfenimide (TBSI) accelerated sulfur-vulcanisation of IIR [26] compared to the TBBS accelerated systems [29]. At low cure times, only A1c and A2c polysulfidic structures (50 ppm) are observed. At longer cure times, A1c and A2c polysulfidic structures reduce in sulfur rank to monosulfide (45 ppm), and B1c (58 ppm). B1t (64 ppm) and C1c (45 ppm) polysulfidic structures are observed. A small amount of cis-to-trans isomerisation was observed, which increased with sulfur content. The reversion reactions of TBSI-accelerated systems result in a lower degree of sulfurisation as opposed to TBBS-accelerated samples. Based on the equilibrium swelling measurements, TBSI is found to be a less efficient accelerator than TBBS. The 13C NMR chemical shifts occurring in the 40 to 70 ppm region of the TBBS-cured IR system were further examined by the DEPT experiments [30]. Based on the DEPT data along with the APT [31] and chemical measurements, the peak assignments for the resonances at 45.0, 50.2, 50.7 and 58.0 ppm in the previous study were confirmed. The peaks at 50.2 and 50.7 ppm are due to polysulfidic A1c and A2c structures, respectively. The monosulfidic counterparts of these structures are expected at approximately 45 ppm. Thus, the 45 ppm resonance contains contributions from A1c and A2c monosulfides as well as C1c polysulfides. According to observations in the chemical network analysis, it is postulated that the peak at 45 ppm is most likely due initially to the A1c and A2c structures, while at longer cure times, it reflects the formation of polysulfidic C1c structures. The peak at 58 ppm has been assigned to a B1c polysulfide. The resonances at 64 ppm, which has been assigned to a B1t polysulfide, could not be detected in the DEPT experiments probably due to the low concentration. A mechanism of the network formation was proposed based on the data of 13C NMR [30] and chemical network analysis. Figure 9.4 illustrates the concentration profile of (A) accelerator intermediates and (B) accelerator, sulfur and by-products during the vulcanisation of IIR, obtained by high pressure liquid chromatography (HPLC) analysis [32]. Figure 9.5 gives the structures and abbreviations used for the accelerators and cure intermediates. Initially, BtSxBt, which is the product of TBBS/sulfur reaction, reacts with rubber to form A1c and A2c-polysulfidic intermediates along with BtSH. For the later extent of cure, the BtSxZnSxBt structure is produced by the consumption of BtSH with Zn/sulfur, which reacts with the rubber to form B1c and B1t-polysulfidic intermediates. In the reversion period, reduction in sulfur rank and the formation of C1c structure occurs via the following schemes: B1c-Sx-B1c + [BtS]2Zn ←→ B1c-S(x-y)-SBt + B1c-SyZnSBt B1c-S(x-y)-SBt + Rubber → B1c-S(x-y)-C1c + BtSH

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Figure 9.4 The concentration profile of (A) accelerator intermediates and (B) accelerator, sulphur and by-products during the vulcanisation of IR, obtained by the HPLC analysis

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Figure 9.5 Structures and abbreviations used for the accelerators and cure intermediates

The network chain density (νe) and the density of the chains between sulfurisations (2[S]chem) for the NR [27] and high-cis-IIR vulcanised with TBBS were determined by equilibrium swelling and solid-state 13C NMR measurements, respectively. The proportion of the intermolecular crosslinks over a whole range of sulfurisation reactions (Ec) was estimated by the comparison of the chemical network chain density νchem) calculated from the νe using Mullins’s approach with the 2[S]chem obtained by NMR. The Ec in the NR system plotted as a function of cure is shown in Figure 9.6. The Ec is zero until the scorch cure (t10; 10% of rheometer torque increase), increases with cure during the curing periods and hits the maximum, 59.8%, at the optimum cure state (t90). The Ec then decreases with time in the overcuring and drops to 23.4% at 120 minute cure.

9.6 Sulfur-vulcanisation of BR Solid-state 13C NMR spectroscopy was used to study accelerated [33] and unaccelerated [34] sulfur-vulcanisation and sulfur-donor (TMTD) [35] vulcanisation of cis-polybutadiene (BR). Olefinic and methylene carbons of the cis-BR repeating unit typically resonate at 129.5 and 27.5 ppm, respectively. The dominant products occurring in the vulcanisation

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Figure 9.6 The Ec in the NR system plotted as a function of cure

with sulfur alone are the new resonances at 33, 38 and 50 ppm. Similar new resonances are observed in the TMTD accelerated sulfur-cure systems. In the system where zinc oxide was formulated in addition to the elemental sulfur and TMTD, simpler spectra are obtained in which the peaks at 38 and 50 ppm are missing. The peak at 33 ppm is assigned to the trans structure of 1,4-BR. An increasing intensity at 33 ppm peak with cure in both sulfur-cured and accelerated sulfur-cured BR postulates the occurrence of cis-to-trans chain isomerisation in these systems. The resonances at 38 and 50 ppm are assigned to cyclic monosulfide and polysulfidic crosslink structures. The expected monosulfidic junctions are not detected in this study possibly due to the low concentration of these species [33]. Doing the experiment at 75.5 MHz, improved resolution attained in the gated high power decoupling (GHPD) spectra of sulfur-vulcanised BR. Figure 9.7 is the magnified spectrum of the high-cis BR cured with 10 phr sulfur at 150 °C for 30 minutes. Twentyfive new peaks are obtained with cure that are labelled in the Figure 9.7. According to the DEPT experiments, the peaks of no. 1 through 7 are due to methine carbons and all other resonances are due to methylene carbons [34]. Similar vulcanisation products with reduced structural modifications are obtained in the sulfur-donor (TMTD) vulcanised BR [35]. By comparing the GHPD and DEPT results of sulfur-vulcanised BR with the results of sulfur-donor vulcanisation, detailed chemical shift assignments are possible where the peaks from inter- and intra-molecular sulfurisations and those from accelerator fragments can be distinguished.

334


Figure 9.7 The magnified spectrum of the high-cis BR cured with 10 phr sulfur at 150 °C for 30 min. Twenty-five new peaks are obtained with cure which are labelled in the figure

Side reactions including cis-to-trans isomerisation and sulfidic cyclisation are observed along with the formation of crosslinks in the BR cured with sulfur alone. In the sulfurdonor vulcanisation of BR, cis-to-trans isomerisation is the predominant feature of the vulcanisation reaction sequence and seems to obey first-order kinetics with respect to the concentration of accelerator. The chemical crosslink densities determined from the NMR were obtained for both sulfurcured and sulfur-donor cured BR by computing the crosslink densities as one-half the sum of all sulfurised structures, which were then compared with the corresponding crosslink densities obtained by the swelling experiments (modified Flory-Rehner equation). The NMR crosslink densities of the vulcanisates are routinely an order of magnitude greater than the swelling crosslink densities in both systems. It was concluded that sulfuror TMTD-vulcanised BR networks must possess a significant proportion of mechanically ineffective vicinal and/or cyclic crosslink structures. Either a free-radical chain reaction mechanism of crosslink induction or the phase-solubility isolation of crosslinking within

335

Spectroscopy of Rubbers and Rubbery Materials discrete localised domains in the rubber would account for the formation of a high percentage of elastically inefficient crosslink structures. The chemistry of the TBBS-accelerated sulfur-vulcanisation of high-vinyl BR was studied by solid-state 13C NMR at 75.5 MHz [36]. The 13C NMR study of the high-vinyl BR (or high-styrene styrene-butadiene rubber (SBR)) is challenging since the regions between 30 and 45 ppm are typically overlapped by the backbone resonances arising from the complex tacticity distribution in the vinyl-vinyl (or vinyl-styrene, styrene-styrene) sequences. Furthermore, the inherent rigid nature of the vinyl-BR (or styrene sequence) structure broadens the resonance lines, which decreases the resolution in the particular region where the sulfurisation products are expected. The effect of the accelerator to sulfur ratio on the vulcanisation chemistry was also investigated by comparing the vulcanisation products from conventional, SEV and EV formulation systems. With increase in accelerator to sulfur ratio (from conventional to EV) there is a lowering in the sulfur rank. Also, the cis-to-trans isomerisation increases with the amount of accelerator. High resolution MAS techniques of 13C, DEPT, correlated spectroscopy (COSY), total correlation spectroscopy (TOCSY), heteronuclear chemical shift correlation (HETCOR) were used to examine conventional CBS and efficient TMTD vulcanisation of polybutadiene [37]. In conventional CBS vulcanisation, the major vulcanisate 13C NMR peak occurred at 44.9 ppm and was assigned to a trans allylic structure (–C=C-C-Sx with X=3 or 4). The efficient TMTD vulcanisation yielded as main product a 13C NMR peak at 54.0 ppm and was assigned to a cis allylic vulcanisate (-C=C-C-Sx; x=1). While cyclic sulfur by-products were observed in both vulcanisation systems, the CBS formulations gave rise to a higher percentage postulated to be formed via a episulfide intermediate. The vulcanisate structure of polyoil 110 after curing with CBS was examined using 1D and 2D NMR experiments of 1H, 13C, COSY, and HETCOR [38]. Polyoil 110 is a butadiene oligomer containing 72% cis, 27% trans, and 1% vinyl structures. Assignments were made in the 1H spectrum using model data and COSY NMR results and then correlated to the 13C spectrum. Assignment in the 13C spectrum showed monosulfidic crosslink peaks at approximately 41 ppm, disulfides between 50-52 ppm, and polysulfides between 52-55 ppm. The author concluded that 13C shift increments obtained from previous model studies were not useful in examining 13C network structure. The vulcanisation of high cis-BR was studied in unfilled and carbon black filled samples vulcanised with TBBS at 148 °C at variable curing times of 0, 20, 30, and 60 minutes [39]. On examining the individual peaks in the 13C spectrum (Figure 9.8), with the onset

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Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR of vulcanisation, two large peaks at 31.4 and 40.9 ppm appear simultaneously, increase and than decrease with time. The two peaks closely integrate to the same values and probably represent a CH2-CHS- pair. From chemical shift addition values reported in the literature [36] the closest assignment would be a monosulfide-linked carbon. This appears an unlikely assignment as monosulfides are not generally formed early in the cure process. Later in the vulcanisation, grouped methine peaks between 48.9-51.8 and 51.8-55.8 ppm appear although with their corresponding b-CH2 groups between 35.839.2 ppm. The 40.9 ppm peak appears at greater concentration earlier and decreases faster in the carbon black CB filled samples. After 60 minutes cure, the concentration in both unfilled and (CB) filled is the same. In the vulcanisate carbon concentrations between 35.8-39.2, 48.9-51.8, and 51.8-55.8 ppm, this same trend is observed with unfilled and filled samples. Cis-trans isomerisation increased by approximately 400% during vulcanisation from the starting (0 minutes) and ending (60 minutes) cure times as determined from the trans peak areas at 33.0 ppm. Vulcanisate structures of BR crosslinked with cyclic disulfides was studied by NMR.36 Using high resolution MAS techniques of DEPT, COSY, TOCSY, and HETCOR, the resulting spectra showed that crosslinking gave an addition product to the double bond and not the allylic structure found in typical sulfur vulcanisations.

Figure 9.8 The 13C NMR spectra of TBBS vulcanised BR at cure times of 0, 20, 30 and 60 minutes

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Spectroscopy of Rubbers and Rubbery Materials NMR studies of blends of NR and BR have also been reported [37, 40]. It was found that peaks characteristic of BR crosslinking appear at early stages in vulcanisation and that NR crosslink peaks appear much later [37]. Cis-trans isomerisation was found to occur in both rubbers but to a greater extent in BR. Interfacial reactions are observed [38].

9.7 Sulfur-vulcanisation of SBR The NMR characterisation of SBR has also been studied although to a lesser extent than NR and BR. One study was carried out on the model compounds of cyclooctene, cis-3-heptene, trans4-octene, cis-1,7-diphenyl-3-heptene, and cis-1-phenyl-3-heptene [38]. The model compounds were vulcanised with 4.5 phr sulfur, 4.5 phr CBS, and 13.6 phr zinc oxide at 150 °C. The BR models of cyclooctene, cis-3-heptene and trans-4-octene gave vulcanisation products via allylic hydrogen substitution yielding both cis and trans structures. The SBR models of cis-1,7-diphenyl-3-heptene and cis-1-phenyl-3-heptene gave no vulcanisation products under these conditions. Only when the more reactive accelerator, TMTD, and a longer vulcanisation time of 1 hour were used did cis-1-phenyl-3-heptene form vulcanisation products consisting exclusively of allylic trans configuration. Thus vulcanisation near the phenyl group is severely restricted. The 13C NMR examination of CBS cured SBR was performed on unfilled, silica filled, and carbon black filled samples [41]. Several different SBR samples with respect to styrene content and cis, trans, and vinyl BR content were used. The unfilled SBR samples gave 3 major peaks that appeared at 51.0, 50.2, and 49.3 in a spectrum similar to BR vulcanisation. Unfortunately, peaks below 45 ppm are obscured by the different main chain structural peaks of SBR. A difference was seen in the rate of formation of these peaks in filled samples with silica inhibit the vulcanisation rate compared to carbon black filled samples.

9.8 Peroxide, Radiation, and High Pressure Vulcanisation Dicumyl peroxide vulcanisation of high-cis BR and NR samples at different peroxide levels were investigated by solid-state 13C NMR [42]. As the peroxide level increased, decreased signal intensity and peak broadening was observed in the main chain peaks in both BR and NR due to decreased segmental motion. In addition, cis-trans isomerisation was observed during the vulcanisation process in both rubbers. New peaks at 35 and 44 ppm were observed in the BR spectra while peaks at 14.9, 21.4, 30.6, 37.5, and 45.0

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Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR were observed in the NR spectra. While no attempt was made to assign the peaks to crosslink structures, the number and position of the peaks indicate that the vulcanisation mechanism is more complicated than a simple combination of allylic radicals. Network structure and reaction mechanisms in high pressure vulcanisation (HPV) and peroxide vulcanisation of BR was studied by 13C solid-state NMR [43]. Different samples of polybutadiene (51% trans, 38% cis, and 11% vinyl) were peroxide cured with dicumyl peroxide on a silica carrier and by the HPV conditions of 250 °C and 293 MPa. The 13C NMR spectra from peroxide and HPV cures were compared to a control samples heated to 250 °C for 6 minutes under atmospheric pressure. Although no new isolated strong peaks were detected in either the peroxide or HPV vulcanisations, small increases in both spectra were observed at 29.5, 36.0, 46.5, and 48.0 ppm. These peaks compare favourably with calculated shifts from structures that arise from main chain radical addition to the pendent vinyl groups. These assignments are further reinforced by the observation that the vinyl carbon concentration is substantially reduced during vulcanisation in both peroxide and HPV curing. Two peaks at 39.5 and 42.5 ppm appear only in the peroxide spectrum. Cis-trans isomerisation was absent in both cures. The networks of dicumyl peroxide cured cis-polyisoprene in the presence and absence of the coagent, triallyl cyanurate, were investigated by solid-state NMR techniques [44]. In peroxide only curing, peak broadening and cis-trans isomerisation increased with increasing peroxide level in the cure formulation. In addition, an increase in the fractions of addition to the double bond is observed with increasing peroxide levels as observed by the rate of decrease of the olefin peaks in the NMR spectra. When the coagent was added, new peaks at 22.8, 36.0, and 47.6 ppm were observed in the 13C spectra and assigned to different structures formed between coagent and polymer backbone. Solid-state 13C NMR was used to characterise the structure produced in gamma irradiated NR [45]. Cis-trans isomerisation increases from 5.1% at 40 Mrad to 10.3% with 161 Mrad. Peaks in the NMR spectra were observed during irradiation due to the formation of a quaternary carbon (44 ppm) and its attached methyl group (21 ppm), a methine carbon (38 ppm), and a g-carbon of a cis isomeric unit adjacent to a trans isomeric unit (30.1 ppm). The mechanism of radiation vulcanisation of NR with 2-ethyl hexyl acrylate (EHA) was examined by 13C NMR [46]. All peaks in spectrum due to the acrylate decrease in intensities and broaden with increasing radiation dose. Using the solid-state CP technique, all peaks corresponding to EHA and NR were observed. The olefin peak at 129 ppm in the acrylate was absent, however, indicating that all the EHA had polymerised.

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9.9 Vulcanisation of Other Elastomer Systems Microstructural changes of an accelerated sulfur vulcanisation of IIR with TMTD/ZnO/ sulfur has been studied by solid-state 13C NMR spectroscopy [47]. The IIR containing 2% isoprene and 98% isobutylene were formulated using EV and cured at 160 oC for several cure times. The resonances at 20.3 and 24.4 ppm, which are due to trans isoprene units in the IIR, decrease with cure, while the resonances at 26.9 and 25.2 ppm which arise from cis isoprene units increase with cure times. The cis:trans ratio increases up to a maximum ratio of approximately 4:1 at a cure time of 60 minutes. New resonances are observed at 15, 21, 23.6 and 49 ppm. The peak at 49 ppm is assigned to the mixture of the isoprene units in cis-IIR, polysulfidic A1t and polysulfidic A1c structures. The resonance peaks at 15, 21 and 23.6 ppm are assigned to the isoprene units in mono- and polysulfidic B1t, mono- and polysulfidic B1c and polysulfidic A1t structures, respectively. No reaction occurs in the isobutylene units. No migration of the double bond saturation, internal cyclisation or sulfurisation resulting in C1t and C1c structures is observed. The curing of brominated poly(isobutylene-co-4-methylstyrene) by pure ZnO was studied by solid-state 1H and 13C MAS NMR [48]. From NMR results, crosslinking appears to occur by aromatic ring addition of the benzylic carbocation generated from an intermediate PhCH2-Br-Zn complex. Solid-state 13C NMR has been used to identify elastomers in binary blends of chloroprene (CR) and NR, CR and CSM, NR and CSM, and SBR and acrylonitrile-butadiene rubber (NBR). The type of NBR can be determined by identifying the sequences of acrylonitrile and butadiene. The tertiary blend of NR/SBR/BR was also studied [49]. High-temperature 13 C solid-state NMR identified ethylene-propylene diene terpolymer (EPDM) and fluoro and nitrile rubbers [50]. Elastomeric components and compositions in BR/SBR and NR/BR/SBR blends have been studied by 13C solid-state NMR. The MAS spectra are of sufficient quality for polymer identification of the carbon black filled vulcanisates in most cases [51]. Vulcanised SBR/EPDM diblends were quantitatively characterised by solid-state 13C NMR spectroscopy. The SBR/EPDM blend ratio can be determined as well as the cis-1,4, trans1,4 and vinyl-1,2 butadienes and styrene ratios in the SBR and the ethylene and propylene contents in the EPDM. No evidence for homo- and co-vulcanisation has been obtained in these systems. No evidences are found for the change in cis:trans ratio in the SBR upon the vulcanisation [52].

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9.10 Effect of Carbon Black on Vulcanisation of Elastomers The incorporation of carbon black into elastomeric systems is a process of significant commercial importance. However, the additional stiffness of the sample imparted by the reinforcement effect of fillers is not favourable in terms of the experimental conditions for high-resolution NMR spectroscopy. Electric conductivity of the carbon black may also interfere to some extent. Under these circumstances, filled formulations are not widely used for the study of elastomer vulcanisations where high resolution and signalto-noise ratios are required to detect small amounts of vulcanisation products. Solid-state 13C NMR spectra of carbon black filled, uncured and sulfur-vulcanised IIR were recorded at 22.6 MHz. The line broadening of the filled polymer relative to the unfilled polymer is attributed to incomplete motional narrowing of the NMR lines [53, 54] Incorporation of filler also results in a decrease in the signal-to-noise ratios in the spectra, but fundamentally it does not obscure the qualitative and quantitative nature of the spectra for the moderately cured elastomer systems. Carbon black filled NR vulcanisates have been studied by high-resolution solid-state 13C NMR at 75.5 MHz. It is shown that the detection of poly- and mono-sulfidic crosslinks with 13C MAS spectroscopy is also possible in the case of carbon black filled NR vulcanisates. Quantitative results suggest that the addition of carbon black leads to a decrease in crosslink density with increasing filler content [55]. The effect of carbon black on the vulcanisation chemistry has been studied for sulfurcured NR systems by 13C solid-state GHPD experiments at 75.5 MHz [27]. Figure 9.9 shows the 13C NMR spectra of carbon black (N347) filled NR cured to t50 (50% of rheometer torque increase)(A) and for 120 minutes (B). Similar new resonances are observed as compared to the results for unfilled cis-polyisoprene systems. The increase in the 16 ppm peak in the overcuring period is attributed to the cis-to-trans isomerisation. Incorporation of carbon black does not affect the level of this isomerisation reaction. The peaks at 51 and 45 ppm are due mainly to polysulfidic A1c and monosulfidic A1c structures. Figure 9.10 illustrates the change in peak intensities at 51 ppm (A) and 45 ppm (B) for four formulations with different carbon black loading as a function of cure. The amount of A1c polysulfide decreases with increase in carbon black (Figure 9.10 (A)), while more monosulfidic A1c structures occur after t90 with higher amounts of filler (Figure 9.10 (B)). The occurrence of the B1-type sulfides increases with the level of carbon black loading. These changes in the quantitative nature of the vulcanisation are partially explained by the preferential adsorption of the rubber molecules on the surface of the carbon black aggregates [27].

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Figure 9.9 The 13C NMR spectra of carbon black (N347) filled NR cured to t50 (50% of rheometer torque increase) (A) and for 120 min (B)

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Figure 9.10 The change in peak intensities at 51 ppm (A) and 45 ppm (B) for four formulations with different carbon black loading as a function of cure

Similar trends have been observed in the carbon black (N347) filled, TBBS accelerated sulfur-vulcanisation of high-cis-IIR. In contrast to the NR/CB system, the reversion reactions, i.e., the cis-to-trans isomerisation and the chain scission at 3,4-isoprene units, increase with black content during the overcuring. The amount of the sulfurisation occurring in the carbon black (N347) filled, TBBS accelerated sulfur-vulcanisation of NR and IR have been determined by quantitative

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Spectroscopy of Rubbers and Rubbery Materials measurements in the solid-state 13C NMR [27]. The density of the chains between sulfurisations (2[S]chem) obtained from the 13C NMR was then compared with the effective network chain density (ne) determined by the equilibrium swelling measurements (modified Flory-Rehner equation). The 2[S]chem constantly increases with cure even in the overcuring, while the ne hits the ceiling after t90. The 2[S]chem are almost constant for all formulations with different content of carbon black, suggesting that the incorporation of carbon black does not affect the total amount of chemical sulfurisation reactions. The physical network chain density (nent) was estimated from the comparison of the NMR and swelling results. A linear relationship of the nent in the NR system as a function of carbon black content exists. This result suggests that the physical adsorption plays a major role in the polymer-filler interactions. Similar NR formulations, but using N110, N220, N326, N330, N550 and N765 carbon blacks, have been studied by solid-state 13C NMR and equilibrium swelling measurements [27]. Changes in the structural parameters (surface area and structure) of carbon black do not affect the amount of sulfurisation reactions. The nent at t90 are obtained for each formulation and plotted as a function of carbon black content, see Figure 9.11. The comparison of the slopes of the plots and the 300% modulus data (ASTM D1765 [56], D3192 [57]) suggests the use of the plot as a method of evaluating filler reinforcement effects.

Figure 9.11 Physical network chain densities of NR cure to t90 with a variety of carbon black as a function of carbon-black content

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Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR The new resonances, which appear with vulcanisation in the presence of carbon black, are qualitatively the same as observed for the unfilled formulation systems, i.e., new resonances at 50.7, 50.2, and 37 ppm were detected at lower cure times. At longer cure times these three resonances disappeared and new resonances were detected at 58, 64, and 45 ppm along with new peaks at 17 and 12 ppm and the growing of the intensity of the peak at 14 ppm. As the amount of carbon black content increased, the latter features were observed at an earlier stage of cure. At low cure times, doublet peaks at 50.7 and 50.2 ppm were detected along with the peak at 37 ppm, and these peaks disappeared at longer cure times. The intensities of these two peaks were comparable for every detected pair. In the previous studies, the peaks at 50.7, 50.2, and 37 ppm were assigned to a quaternary carbon of A1c polysulfide, a methine carbon of A1c polysulfide, and a β methylene carbon of A1c polysulfide, respectively. While at longer cure times, resonances at 58, 64, 45, 17, and 12 ppm were detected along with increasing intensity of the 14 ppm peak. The peaks at 58, 64, 17 and 12, and 14 ppm are assigned to B1t monsulfide (α methine), B1t polysulfide (α methine), B1c polysulfide (β methylene), and B1t polysulfide (β methylene), respectively. Accordingly, two phenomena are observed involving the effect of carbon black. One is the reduction in sulfur rank for the A1c sulfide structure, and the other, is the formation of the A1 type product at an earlier cure stage and an increase in the B1 products at a later stage. The reduction in sulfur rank was interpreted as due to an entropy change. The disulfide crosslinks have less mobility than the monosulfide crosslinks due to the rigid nature of -S-S- linkage. Similarly, the mobilities of the polysulfidic crosslinks are probably lower than that of the monosulfidic crosslinks. Therefore the polysulfidic crosslink would occupy more space in the network structure than the monosulfidic crosslink. Considering this fact, the change in sulfur rank from polysulfide to monosulfide would occur along with the increase in entropy when the curing process is extended. Since the system becomes more rigid by the incorporation of the carbon black, the reduction in sulfur rank to monosulfide would eventually be accelerated. The second phenomenon due to carbon black on a proposed scheme based on the vulcanisation process shown in Figure 9.12. During the vulcanisation process, the reactions (1) and (2), and (3) and (4) form polysulfidic A1c and B1t, respectively. Considering the fact that A1c polysulfide reduces its sulfur rank to monosulfide and reaction (3) and (4) is the predominant process in the later stage of cure, the system may favour reaction (3) and (4) rather than (1) and (2) in the whole reaction scheme. BtSH is formed from the

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Spectroscopy of Rubbers and Rubbery Materials reaction producing polysulfidic A1c at the earlier stage of cure, which is the reactant for reaction (4) producing the B1 structure. As previously mentioned, the sulfur and accelerator to rubber ratio may be higher for the system with high carbon black content. As a result, reactions (1) and (2) would occur more rapidly and, consequently, produce enough BtSH for the initiation of reactions (3) and (4) , which means that the B1t polysulfides initiate at the earlier stage of cure for the sample with higher filler content. Since reactions (3) and (4) would occur until the elemental sulfur is consumed, the total amount of B1 structure would be larger for the sample with higher carbon black content. The network chain density measured by equilibrium swelling includes both the chemical and physical crosslinks while the 13C NMR measurements determine only the chemical crosslinks. The observed differences in network chain density between the results of the two measurements (νSWELL - νNMR) reflect differences in the physical crosslinks induced by entanglements involving polymer-polymer interactions and polymer-carbon black interactions. The magnitude of the differences between the two methods of measuring the network chain density increases with the volume fraction of the carbon black for the same stage of cure. The differences in network density values reach a maximum near the 90% cure stage for the filled formulations. These results coincide to similar observations from tensile tests where the maximum filler reinforcement effect occurs at an optimum

Figure 9.12 Mechanism of vulcanisation

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Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR (90%) stage of cure. Consequently, the comparisons of the network densities determined by the swelling and 13C NMR measurements give a measure of the effectiveness of the filler reinforcement at the molecular level.

9.11 Effect of Silica Filler on Vulcanisation Chemistry Solid state 13C NMR spectroscopy was used to identify and characterise the effects that silica has on the network structures that form during the vulcanisation of zinc activated, sulfur-cured cis-1,4 IIR [58]. It was determined that cis-trans isomerism and chain scission were enhanced when silica was added as a filler. The monosulfidic linkages, especially the A1cis and the B1trans type monosulfide linkages, increased in concentration in the silica-filled systems. Also, the polysulfidic linkages, namely the A1cis and B1trans type linkages, decreased in concentration. Finally, it was also determined that the intermolecular crosslinks decreased in the presence of silica, while the intramolecular structures, including pendent side groups and cyclic sulfides, increased in the presence of silica, resulting in a lower overall crosslink density when silica is used as a filler. When silica is used as a filler, the cure ingredients are adsorbed, causing the initial overall lower cure rate. Several ways exist to counteract this effect including inhibiting the silica absorption by adding a silane coupling agent, or polyethylene glycol (PEG). Both additives have been shown in the literature to improve the vulcanisate properties. Therefore, solid state 13C NMR spectroscopy was used to identify and characterise the effects on the network formation when additives, such as a coupling agent and PEG, were added along with silica filler [59]. It was determined that cis-trans isomerisation was inhibited by the inclusion of both the coupling agent and the PEG. However, there was only a slight influence on the chain scission of the main chain carbons. The presence of coupling agent or PEG both caused an increase in the concentration of the polysulfidic structures, especially the A1cis and B1trans type linkages, as well as a decrease in the monosulfidic structures, namely the A1cis and B1trans type structures. Therefore, the intermolecular crosslinks, mainly the polysulfidic structures, increased with the use of the additives, while the intramolecular structures, which are mainly monosulfidic in nature, decreased. Therefore, the network chain density of the silica-filled rubber system increased when additives were included in the formulation.

9.12 Thermal-Oxidation of Network Structures Elastomeric materials undergo both thermal and oxidation degradation over time. Main chain scission and loss of sulfur crosslinks can occur with either factor or by both factors by a thermo-oxidative mechanism.

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Spectroscopy of Rubbers and Rubbery Materials The chemical changes that occur in thermal and oxidative degradation of sulfur cured NR were investigated using 13C NMR analysis [58]. The TBBS vulcanised samples of NR were heated at 100 °C and 150 °C under both air (oxidative) and nitrogen (inert) environments to assess both temperature and oxidative effects. Changes were observed by measuring peak areas under the crosslinked carbon resonance peaks. A and B-type monosulfide linkages were found to decompose much faster by oxidative factors at both 100 °C and 150 °C than under nitrogen at either temperature. The polysulfide linkages, however, showed the opposite result of thermal degradation predominating over oxidative factors. Cis-trans isomerisation of the main chain carbons was observed at the higher temperature while absent at 100 °C. Another study used 13C NMR to examine the thermal ageing of NR with both peroxide and conventional and EV sulfur vulcanised samples [59]. The samples were heat aged at 70 °C in air with 13C NMR recorded after 3 and 21 days. The vulcanisate structures vulcanised by peroxide and the EV sulfur system showed low sensitivity to ageing effects while peaks in the conventional sulfur system almost disappeared after 21 days of ageing. In addition, the main chain isoprene carbon peaks showed considerable broadening due to increasing stiffness of the network.

9.13 Summary From this chapter, it is clear that high resolution 13C NMR spectroscopy has made a large number of contributions to our knowledge of the structure of vulcanised crosslinked elastomers as well as the mechanism by which the vulcanisation process occurs. It is anticipated that further NMR measurements of these systems will continue to generate new structural and mechanistic information.

Acknowledgements The authors wish to acknowledge the support of the National Science Foundation as well as the support of the Yokohoma Rubber Company and the Pirelli Rubber Company.

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10. M. Mori and J.L. Koenig, Annual Reports in NMR Spectroscopy, Academic Press Ltd., Vol. 34 (1997) 11. M. Andreis and J.L. Koenig, Advances in Polymer Science, 1989, 89, 71. 12. R.A. Kinsey, Rubber Chemistry Technology, 1990, 63, 407. 13. M.R. Krejsa and J.L. Koenig, Rubber Chemistry and Technology, 1993, 66, 376. 14. M.R. Krejsa and J.L. Koenig, Elastomer Technology Handbook, CRC Press, 1993, Chapter 11. 15. J.L. Koenig and D.J. Patterson, Elastomers Rubber Technology, 1987, 32, 31. 16. A.M. Zaper and J.L. Koenig, Rubber Chemistry and Technology, 1987, 60, 252. 17. The Sadtler Guide to Carbon-13 Spectra, Ed., W.W. Simons, Sadtler Research Laboratories, 1983. 18. W. Gronski, H. Hasenhindl, B. Freund and S. Wolff, 1991, Kautschuk und Gummi Kunststoffe, 44, 119. 19. M. Andreis, J. Liu and J.L. Koenig, Journal of Polymer Science, Phys. Ed., 1989, 27, 1389. 20. Natural Rubber Science and Technology, Ed., A.D. Roberts, Oxford University Press, New York, 1988, Chapters 11-12. 349

Spectroscopy of Rubbers and Rubbery Materials 21. C.G. Moore, L. Mullins and P.M. Swift, 1961, Journal of Applied Polymer Science, 5, 293. 22. A.Y. Coran, Rubber Chemistry Technology, 1995, 68, 351. 23. A.Y. Coran, Science and Technology of Rubber, Academic Press, 1994, Chapter 7. 24. A.M. Zaper and J.L. Koenig, Rubber Chemistry and Technology, 1987, 60, 278. 25. D. Parker and J.L. Koenig, Journal of Applied Polymer Science, 1998, 70, 1371. 26. M.R. Krejsa and J.L. Koenig, Rubber Chemistry and Technology, 1993, 66, 73. 27. M. Mori and J.L. Koenig, Journal of Applied Polymer Science, 1998, 70, 1391. 28. W. Gronski, U. Hoffmann, G. Simon, A. Wutzler and E. Straube, Rubber Chemistry and Technology, 1992, 65, 63. 29. M.R. Krejsa and J.L. Koenig, Rubber Chemistry and Technology, 1992, 65, 427. 30. M.R. Krejsa and J.L. Koenig, Rubber Chemistry and Technology, 1994, 67, 348. 31. R.C. Hirst, ACS Rubber Division Meeting, October 1991, Paper No. 69. 32. A.B. Sullivan, C.J. Hahn and G. H. Huhls, Rubber Chemistry and Technology, 1992, 65, 488. 33. A.M. Zaper and J.L. Koenig, Macromolecular Chemistry, 1988, 189, 1239. 34. R.S. Clough and J.L. Koenig, Rubber Chemistry Technology, 1989, 62, 908. 35. S.R. Smith and J.L. Koenig, Rubber Chemistry and Technology, 1992, 65, 176. 36. M.A. Rana and J.L. Koenig, Rubber Chemistry and Technology, 1993, 66, 242. 37. R. Hulst, R.M. Seyger, J.P.M. van Duynhoven, L. van der Does, J.W.M. Noordermeer, and A. Bantjes, Macromolecules, 1999, 32, 22, 7521. 38. J. Hahn, M. Runk, M. Schollmeyer, U. Theimer and E. Walter, Kautschuk und Gummi Kunststoffe, 1998, 51, 206. 39. Dallas D. Parker, Makio Mori and J.L. Koenig, submitted to Rubber Chemistry Technology,

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Chemical Characterisation of Vulcanisates by High-Resolution Solid State NMR 40. B. Klei and J.L. Koenig, Rubber Chemistry and Technology, 1997, 70, 231. 41. L. Pellicioli, S.K. Mowdood, F. Negroni, D.D. Parker and J.L. Koenig, submitted to Rubber Chemistry and Technology, 42. D.J. Patterson, J.L. Koenig and J.R. Shelton, Rubber Chemistry and Technology, 1983, 56, 5, 971. 43. M. Bellander, B. Stenberg and S. Persson, Journal of Applied Polymer Science, 1999, 73, 14, 2799. 44. S.J. Oh and J.L. Koenig, Journal of Polymer Science, Part B: Polymer Physics, 2000, 38, 11, 1417. 45. D.J. Patterson and J.L. Koenig, Applied Spectroscopy, 1987, 41, 3, 441. 46. D.J.T. Hill, J.H. O’Donnell, M.C.S. Perera, P.J. Pomery and P. Smetsers, Journal of Applied Polymer Science, 1995, 57, 10, 1155. 47. M.R. Krejsa and J.L. Koenig, Rubber Chemistry and Technology, 1991, 64, 40. 48. R.R. Eckman, I.J. Gardner and H.-C. Wang, Rubber Chemistry and Technology, 1993, 66, 1, 109. 49. D. Gross and J. Kelm, Kautschuk and Gummi Kunststoffe, 1987, 40, 13. 50. D. Gross and J. Kelm, Kautschuk und Gummi Kunststoffe, 1985, 38, 1089. 51. R.A. Komoroski, Rubber Chemistry and Technology, 1983, 56, 959. 52. G.P.M. Van Del Velden and J. Lelm, Rubber Chemistry and Technology, 1990, 63, 215. 53. Schaefer, Macromolecules, 1972, 5, 427. 54. J. Schaefer, S.H. Chin and S.I. Weissman, Macromolecules, 1972, 5, 798. 55. M. Mori and J.L. Koenig, Rubber Chemistry and Technology, 1997, 70, 671. 56. ASTM D1765-01, Standard Classification Systems for Carbon Blacks used in Rubber Products, 2001. 57. ASTM D3192-00, Standard Test Methods for Carbon Black Evaluation in NR (Natural Rubber, 2000.

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Spectroscopy of Rubbers and Rubbery Materials 58. C. Hill and J.L Koenig, Polymer Bulletin, 1998, 40, 275. 59. M.L. Kralevich and J.L. Koenig, Composite Interfaces, 1997, 5, 125. 60. J.Y. Buzare, G. Silly, J. Emery, G. Boccaccio and E. Rouault, European Polymer Journal, 2001, 37, 1, 85.

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Characterisation of Chemical and Physical Networks in Rubbery Materials Using …

10

Characterisation of Chemical and Physical Networks in Rubbery Materials Using Proton NMR Magnetisation Relaxation Victor M. Litvinov

Abstract The use of solid-state nuclear magnetic resonance (NMR) transverse magnetisation relaxation experiments for characterising various types of viscoelastic materials and rubbery phases in blends is reviewed in this chapter. The methods applied for quantitative analysis of the density of chemical crosslinks, temporary and trapped chain entanglements and physical network junctions, which are formed in filled rubbers, semi-crystalline and ionic containing elastomers, are discussed. Several examples of the determination of the molecular-scale heterogeneity of polymer networks are given. Use of NMR magnetisation relaxation methods in real-time NMR experiments and quality control are discussed.

10.1 Introduction The density of chemical crosslinks and physical network junctions largely affects the mechanical properties of rubbery materials and polymer blends containing rubbery components. Despite the apparent simplicity, the network structure has a complex topology, which can significantly affect functional properties. The type of curing may often cause significant differences in network topology and properties of cured materials. Different types of network heterogeneity will be found in cured materials if no precautions are taken to control the curing chemistry and conditions. The following types of molecularscale heterogeneity may occur: heterogeneity in the distribution of network junctions, polymer chains unattached to the network, and/or dangling chains and loops. In addition, the type of network junction, i.e., the functionality and bulkiness, which determine the network junction’s ability to fluctuate (affine versus phantom networks), may also affect the mechanical properties. A difference in curing conditions in the sample volume, for example a difference in temperature or in the concentration of vulcanisation agents, will result in spatial heterogeneity of the network structure. Besides chemical crosslinks the following types of physical network junctions may occur in rubbery materials:

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Spectroscopy of Rubbers and Rubbery Materials 1) Temporary and trapped chain entanglements, 2) Junctions that are formed due to chain adsorption at the surface of active fillers, 3) Junctions that are formed by crystallites, 4) Junctions deriving from strong hydrogen and ionic bonds, and 5) Junctions that are formed at the interface of polymer blends and those in materials that reveal nano-scale phase separation. These are variables in the network structure that can be utilised to modify the properties of cured materials. However, they also cause difficulties in the analysis of network structures and complicate efforts to determine structure-property relationships. To be able to relate mechanical properties to the composition of rubbery compounds and curing chemistry it is essential to understand the network topology of the resultant cured materials. The methods that are used to analyse network structures can be generally subdivided into three categories on the basis of their methodology: 1) Characterisation of physical properties of cured materials in relation to volume-average network density. The most common methods are equilibrium swelling and mechanical measurements [1-6]. Besides information on the network structure, the mechanical methods provide information that is useful for practical applications, such as information on the modulus of elasticity, ultimate tensile properties and glass transition temperature (Tg). Rubber elasticity theory and phenomenological theories are used to relate a measured quantity to the density of chemical and physical crosslinks. Several molecular models have been developed for ‘ideal’ defect-free networks [1, 7-9]. The validity and applicability of these models is however a heavily debated topic and several discrepancies between theory and experiment still remain, e.g., the role of chain entanglements, network defects and network heterogeneity. It is generally acknowledged that traditional methods are not capable of providing fully reliable information on network topology [9, 10, 11]. 2) Analysis of chemical conversion and cure chemistry is another way of studying network structures. Several techniques are used for this purpose, e.g., optical spectroscopy [12], high-resolution NMR spectroscopy and titration of non-reacted functional groups. The spectroscopic methods can be used for quantitative analysis of crosslinks [13-15]. Chemical conversion is usually closely related to the network density. However, no exact quantitative information on the network structure can be obtained because reacted groups can form

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Characterisation of Chemical and Physical Networks in Rubbery Materials Using … not only chemical crosslinks but also viscoelastically ineffective chains, such as chain branches and chain loops. Furthermore, side-reactions, which may readily lead to the formation of additional crosslinks, complicate data interpretation due to the overlapping of signals from different types of chemical groups in complex mixtures. Moreover, physical network junctions are virtually not detected by spectroscopic methods. 3) Analysis of molecular mobility of polymer chains is used to study network density and its heterogeneity. Dynamic mechanical analysis (DMA) and dielectric spectroscopy provide information on the mobility of polymer chains, which is linked to the network density [16]. One of the most informative and sensitive methods of network analysis is solid-state NMR [17-23]. Furthermore NMR imaging or microscopy is used to determine the spatial heterogeneity of rubbery materials on a scale of 15-50 micrometres [19,24-26]. Different types of NMR magnetisation relaxation experiments are used to analyse local and long-range spatial mobility of polymer chains. Since chain motion is closely coupled to the length of network chains, chemical information on network structure and network defects can be obtained in this way. The mobility of polymer chain units of different chemical origins can be determined by means of selective NMR magnetisation relaxation experiments using high-resolution solid-state NMR techniques. These techniques can also be used for the analysis of the network structure in rubbery blends. Relationships between NMR relaxation parameters, dielectric and mechanical properties have been established [27-29]. Evidently, the most comprehensive information on network structure in relation to properties can be obtained by using these three complementary methodologies. The use of solid-state NMR magnetisation relaxation experiments to characterise network structures in various rubbery materials is reviewed in this chapter. Comprehensive reviews of high-resolution NMR techniques can be found elsewhere [21-23, 30-35].

10.2 Network Structure Analysis by Means of NMR Transverse Magnetisation Relaxation Solid-state NMR magnetisation relaxation experiments provide a good method for the analysis of network structures. In the past two decades considerable progress has been made in the field of elastomer characterisation using transverse or spin-spin (T2) relaxation data [36-42]. The principle of the use of such relaxation experiments is based on the high sensitivity of the relaxation process to chain dynamics involving large spatial-scale chain motion in elastomers at temperatures well above the Tg and in swollen networks. Since chain motion is closely coupled to elastomer structure, chemical information can also be obtained in this way.

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Spectroscopy of Rubbers and Rubbery Materials The temperature dependence of the 1H T2 relaxation time of well-defined end-linked poly(propylene oxide) (PPO) networks [43] is shown in Figure 10.1 [44] 1. The distinguishing feature of T2 relaxation in the case of viscoelastic networks is the high temperature plateau that is observed at temperatures well above the Tg [45-49]. At these temperatures, the temperature independence of T2 is attributable to constraints limiting the number of possible conformations of a network chain to those of a free chain. The slight increase in T2 with an increasing temperature, which is observed for some networks in the plateau temperature range, may be attributable to the presence of network defects, such as dangling chains, long chain loops and residual non-vulcanised material [49]. A value of T2 at the plateau, T2pl, is determined by the asymmetry of random rotations of monomer units, and does not depend on the mechanism and frequency of motions of network chains. The theory of the transverse relaxation in elastomeric networks relates T2pl to the number of statistical segments, Z, between chemical and physical network junctions [45, 46]: Z = (T2pl)/[a(T2rl)] where a is the theoretical coefficient, which depends on the angle between the segment axis and the internuclear vector for the nearest nuclear spins at the main chains. In the case of polymers containing aliphatic protons in the main chain the coefficient a is close to 6.2 ± 0.7 [46]. T2rl is the relaxation time measured at temperatures below the Tg of the swollen polymer. A solvent free of hydrogen atoms (fully deuterated, chlorinated or fluorinated solvent) is used in this experiment. The weight average molecular mass of network chains between chemical and physical junctions, <Mc+p>w, can be calculated using the number of backbone bonds in the statistical segment, C∞: <Mc+p>w = ZC∞Mu/n where Mu is the molecular mass per monomer unit and n is the number of rotatable backbone bonds in the monomer units. For copolymers, the average molecular mass per rotatable backbone bond is used in the Mu calculation. The maximum relative error of this NMR network density determination was estimated to be about 15%–25% [49, 50]. Comparative analysis of the network density measured in the same samples with the aid of the NMR method and other techniques, such as stress-

1

Tg for a linear, high-molecular-mass PPO, which was prepared from a polypropylene glycol precursor (with a molecular mass of 4000 g/mol) using a chain extender with a chemical structure similar to that of the crosslinker, is –62.5 °C. The Tg was determined from DMA measurements performed at 0.215 Hz [43].

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Figure 10.1 Temperature dependence of the 1H T2 relaxation time of well-defined endlinked (PPO) networks with narrow molecular mass distributions between chemical crosslinks [44]. The molecular mass of network chains (in g/mol) is shown in this figure. The temperature dependence of a linear, high-molecular-mass poly(propylene oxide) prepared from a polypropylene glycol precursor (with a molecular mass of 4000 g/mol) using a chain extender with a chemical structure similar to that of the crosslinker is shown for comparison. The synthesis of the model networks has been described elsewhere [43]

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Spectroscopy of Rubbers and Rubbery Materials strain measurements, equilibrium swelling and 13C NMR spectroscopy, proves that T2 relaxation data provide quantitative data on the network density [15, 49-52]. As an example, the results obtained for UV-cured acrylates with a broad crosslink density range were compared using the NMR method and DMA (see Figures 10.2 and 10.3 [52]). A fairly good correlation between the storage modulus and the relaxation rate (1/T2) was observed. Both methods gave similar values for the mean molecular mass of network

Figure 10.2 The relaxation rate (1/T2s)max measured for a cured mixture of a poly(ethylene glycol) diacrylate (Mn = 700 g/mol) and 2-ethylhexyl acrylate as a function of the storage modulus at 273 K (-0.1 °C) [52]. The rubbery plateau was observed for all samples at 273 K (-0.1 °C). (1/T2s)max corresponds to the relaxation component with short decay time that was measured at 323 K (50 °C) for partially swollen in 1,1,2,2-C2D2Cl4 samples. This relaxation component corresponds to the relaxation of network chains. The line represents the result of a linear regression analysis: intercept = 1.1 ± 0.3 ms-1; slope = 0.34 ± 0.02 ms-1(MPa)-1. The correlation coefficient equals 0.992 Reprinted with permission from V.M. Litvinov and A.A. Dias, Macromolecules, 2001, 34, 12, 4051, Figure 10. Copyright 2001 American Chemical Society

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Characterisation of Chemical and Physical Networks in Rubbery Materials Using … chains at small values of Mc+e. A substantial difference in the Mc+e value at low crosslink density in caused by the effect of network defects, which decrease volume average network density as determined by DMA. The effect of network imperfections was excluded from the NMR analysis, as will be discussed below (see Section 10.3).

Figure 10.3 Mean molecular mass between chemical crosslinks and trapped chain entanglements Mc+e in a cured mixture of a poly(ethylene glycol) diacrylate (PEGDA) and 2-ethylhexyl acrylate (EHA) as a function of the EHA content [52]. Mc+e values were determined from (1/T2s)max and the plateau modulus (see Figure 10.2). A substantial difference in Mc+e value, as determined by these two methods at low crosslink density, is caused by the effect of network defects which decrease volume average network density determined by DMA (see Section 10.3). The molecular mass of PEGDA (Mn = 700 g/mol) is indicated by an arrow. The molecular mass of network chains in cured PEGDA is about three times smaller than that of the initial monomer. The molecular origin of this difference is discussed in Section 10.3 Reprinted with permission from V.M. Litvinov and A.A. Dias, Macromolecules, 2001, 34, 12, 4051, Figure 11. Copyright 2001 American Chemical Society

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Spectroscopy of Rubbers and Rubbery Materials The effect of chain entanglements on T2 was studied using series of ethylene-propylenediene (EPDM) vulcanisates with a broad range of chemical crosslink densities [49]. The network density in the vulcanisates was determined with the aid of the following methods: 1 H NMR transverse magnetisation relaxation, Mooney-Rivlin analysis of stress-strain curves [2-4] and equilibrium swelling experiments. It was shown that in the case of the original vulcanisates, the NMR method measures the total network density, which comprises chemical crosslinks (CC) and temporary entanglements (EN/TE) and trapped chain entanglements (EN/TR). For partially swollen samples, the NMR method determines the network density comprising CC and EN/TR. A difference in the T2 values obtained for the original vulcanisates and the swollen ones was analysed with respect to entanglement density. The estimated molecular mass of EPDM chains between apparent chain entanglements (EN) was 1900 ± 200 g/mol. The value obtained is in good agreement with the values obtained using other techniques [49]. The CC density was determined from the total density of network junctions and the measured entanglement density on the assumption that CC and chain entanglements are decoupled and additive. The contribution of CC, EN/TE and EN/TR to the total network density was estimated for EPDM vulcanisates with a broad crosslink density range. The results of the T2 relaxation studies prove that this method is a very useful technique for the quantitative characterisation of network structures, while the more sophisticated NMR techniques, which also determine the residual dipole-dipole interactions [31, 53-60], provide specific information for the chemical structure and molecular mobility, which may be useful in determining mechanisms of molecular motions and refining interpretations of the non-selective T 2 relaxation method, especially for composite materials.

10.3 Characterisation of Network Heterogeneity and Network Defects Mechanical properties of crosslinked elastomers are influenced not only by the volumeaverage crosslink density but also by network heterogeneity. The influence of structural defects (such as residual sol, dangling chains, chain loops and the heterogeneity of the junction distribution) on the viscoelastic properties and the equilibrium swelling data is still under discussion. Local methods which probe molecular properties are very suitable for the determination of the degree of network heterogeneity [11]. Several types of heterogeneity may occur in rubbery materials: 1) molecular-scale heterogeneity, which is caused by the chemical heterogeneity of uncured elastomers, network defects and heterogeneous distribution of network junctions on a molecular level;

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Characterisation of Chemical and Physical Networks in Rubbery Materials Using … 2) morphological heterogeneity of rubbery compounds due to a spatially heterogeneous distribution of components and filler in the compound; 3) spatial heterogeneity due to differences in curing conditions such as temperature and concentration of vulcanisation agents throughout the sample volume. A significant difference between the large spatial-scale mobility of network chains and that of network defects allows us to determine the degree of network heterogeneity. The most reliable data are obtained for swollen samples, because an increasing solvent content results in the disentanglement of network defects from network chains [52, 61, 62]. The molecular mobility of network chains is consequently decoupled from that of network defects, resulting in a major distinction in the relaxation behaviour. The effect of a solvent on the T2 relaxation decay of cured acrylates is shown in Figure 10.4 [52]. The monoexponential T2 decay of a cured acrylate is split into two components upon swelling. This behaviour is most clearly observable at high solvent concentrations [52, 62]. One component has short decay time (T2s), which is comparable with that of the original sample. This component apparently derives from network chains. Starting at a low volume solvent content Vs, (see Figure 10.5), T2s shows an increase, which may be attributable to the following phenomena: 1) the disentanglement of network chains, 2) an increase in the frequency of the large spatial-scale chain motion, and 3) a slight decrease in the strength of the inter-chain proton dipole-dipole interactions. At Vs ≈ 40 - 50 vol.%, T2s reaches a maximum value. A value of T2s at the maximum, (T2s)max, is related to the molecular mass of the network chains between chemical crosslinks and trapped chain entanglements [49]. At higher values of Vs, T2s decreases until the state of equilibrium swelling is reached. This decrease in T2s is thought to reflect the increase in the inter-chain proton dipole-dipole interactions as a result of the network chain elongation following a progressive increase in the solvent fraction in a swollen gel [20]. The long decay time (T2l) of the other component is typical of semi-diluted polymer solutions. The T2l value continuously increases with an increasing solvent content. This

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Figure 10.4a

Figure 10.4b

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Figure 10.4 1H T2 decay measured for a cured mixture of a poly(ethylene glycol) diacrylate and 2-ethylhexyl acrylate (10:90 wt%) without a solvent (a) and with 46 vol.% (b) and 90 vol.% (c) 1,1,2,2-C2D2Cl4 [52]. The solid line represents the result of a least-squares adjustment of the decay with a liner combination of two exponential functions. The dotted lines represent separate components Reprinted with permission from V.M. Litvinov and A.A. Dias, Macromolecules, 2001, 34, 12, 4051, Figure 6. Copyright 2001 American Chemical Society

component apparently originated from the relaxation of network defects, which are disentangled from network chains in a swollen state. At the equilibrium swelling degree, the relative fraction of the T2l relaxation component could be used as a measure of the fraction of highly mobile network defects. The described behaviour of a T2 relaxation decay, following a progressive increase in the solvent fraction, is typical of networks containing a significant fraction of network defects [52, 62]. However, precautions should be taken in the event of such analyses of T2 relaxation with respect to network heterogeneity, as will be discussed later in this section. Distinct T2 relaxation components with widely differing mean decay times suggest molecular or macroscopic heterogeneity of the material. In such cases the submolecule concept can be used to describe the relaxation behaviour [20]. In a simplified interpretation, the overall T2 relaxation decay of a heterogeneous elastomer is the weighted sum of the decays originating from the submolecules, which are defined as the network

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Figure 10.5 The T2 relaxation time for a polymer network with defects against the volume fraction of a good solvent

chains that are formed by the chemical and physical junctions, and network defects, i.e., chains that are not attached to the network, chain loops and dangling chain ends. The large spatial-scale mobility of these submolecules differs substantially, and so does their relaxation behaviour. The relative contribution of the submolecules to the overall decay is proportional to the number of protons attached to these chain fragments. A quantitative analysis of the shape of the decay curve is not always straightforward due to the complex origin of the relaxation function itself [20, 36, 63-66] and the structural heterogeneity of the long chain molecules. Nevertheless, several examples of the detection of structural heterogeneity by T2 experiments have been published, for example the analysis of the gel/sol content in cured [65, 67] and filled elastomers [61, 62], the estimation of the fraction of chain-end blocks in linear and network elastomers [66, 68, 69], and the determination of a distribution function for the molecular mass of network chains in crosslinked elastomers [70, 71]. The applicability of T2 experiments for determining the molecular mass distribution of network chains is illustrated in Figure 10.6.

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Figure 10.6 The distribution function of the chain length for two EPDM sulfur vulcanisates whose crosslink densities differ by a factor of about two and that for the physical mixture (50:50 mass%) of these vulcanisates [72]

Figure 10.6 shows the distribution function of the chain length for two EPDM sulfur vulcanisates whose crosslink densities differ by a factor of about two and the distribution factor for the physical mixture (50:50 mass %) of these vulcanisates [72]. The vulcanisates show a rather broad distribution in molecular mass of network chains, which is apparently caused by the statistical distribution of the third monomer in the original EPDM and possibly by incomplete conversion of the third monomer’s double bonds during vulcanisation. The physical mixture of the vulcanisates shows a bimodal distribution function with maxima at about the same chain length as that for the two vulcanisates in the mixture. This shows that the method is suitable for analysing heterogeneous networks. The analysis of computer-simulated T2 decay curves of bimodal networks shows that a difference in the molecular masses of network chains greater than a factor of two can be determined with the aid of this method. Bulky crosslinks or side-groups in the network chains, e.g., dendritic wedges [73], may also influence molecular mobility and viscoelastic properties of polymer networks. For example, UV curing of difunctional acrylates results in the formation of zip-like network junctions, which may be regarded as extreme cases of bimodal networks [52]. Results obtained with the NMR T2 relaxation method agree well with those of mechanical tests

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Spectroscopy of Rubbers and Rubbery Materials for molecular mass between crosslinks, Mc+e, as was discussed in Section 10.2 (see Figure 10.3) [52]. However, the values of Mc+e obtained are significantly lower than the molecular mass of the initial diacrylate. So, network junctions of this type evidently strongly affect elastic properties and T2 relaxation. As anticipated, classical rubber elasticity theories cannot be used to characterise heterogeneous networks.

10.4 Network Structure in Oil-Extended Rubbers - Effect of Chain Entanglements Knowledge of the physical state of oil in oil-extended rubbers and their vulcanisates is of great importance, because the processing of oil-extended rubbers and the mechanical properties of such rubbers are strongly affected by the rubbers’ oil content and the type of oil concerned. 1 H T2 relaxation of oil-extended EPDM revealed two distinct relaxation components whose characteristic decay times are comparable with those of initial rubber and paraffinic oil (Figure 10.7) [74]. This suggests that the components with a short and long decay time mainly originate from the relaxation of rubbery chains and oil molecules, respectively. Up to 400 K (127 °C), the fraction of the component with a long decay time was smaller than the concentration of oil hydrogen in the oil-extended rubber. Apparently, a small fraction of oil molecules shows a molecular mobility comparable with that of EPDM chains. The fraction of these physically trapped oil molecules decreased with an increasing temperature, and

Figure 10.7 The T2 relaxation decay of oil-extended EPDM rubber containing 12.2 (o), 33.7 (❑) and 50.0 mass% (Δ) paraffinic oil [74]. The line corresponds to the best fit of experimental data points with a linear combination of two exponential functions

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Characterisation of Chemical and Physical Networks in Rubbery Materials Using … above 400 K (127 °C) the molecular mobilities of the EPDM chains and oil molecules were largely decoupled. It was shown that 1H T2 relaxation experiments above this temperature can be used to quantitatively determine both the oil content of EPDM rubbers and the molecular mass of network chains in oil-extended vulcanisates. The results obtained suggest that the molecular diffusion of oil molecules depends largely on: 1) the mesh size of network holes with respect to the size and shape of oil molecules, 2) the molecular mobility of EPDM chains, and 3) intermolecular forces of cohesion. The network densities of the oil-extended EPDM rubbers were determined by means of 1 H T2 relaxation experiments [74]. The EPDM samples containing varying amounts of paraffinic oil (from 5 to 100 phr) were cured under the same conditions, using the same vulcanisation recipe. The network density was nevertheless found to decrease substantially with an increasing oil concentration in the vulcanisates (Figure 10.8). Analysis of these data using scaling laws for polymer melts and solutions [75-77] suggests that the decrease in the overall network density, comprising chemical crosslinks and chain entanglements, is due mainly to a decrease in the entanglement density with an increasing oil content.

Figure 10.8 The density of network junctions as a function of the volume fraction of paraffinic oil in EPDM/oil vulcanisates [74]. The solid line represents the result of a linear regression analysis of the dependence: (intercept = 453 ± 5 mmol/kg; slope = -6.2 ± 0.0.3 mmol/kg; the correlation coefficient = 0.996). Maximum torque in the rheometer curve for the vulcanisates is shown on the right ordinate

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10.5 Network Structure in Filled Rubbers - Rubber-Filler Interface and the Structure of the Physical Network

10.5.1 NMR Relaxation of Filled Rubbers The incorporation of fillers in rubbers is of significant commercial importance, since use of fillers not only enhances the end product’s mechanical properties but can also decreases its costs. Although much research has been carried out on filled rubbers using different techniques, the molecular origin of the filler’s reinforcing effect is still unclear. The reinforcing effect at a moderate strain (, has been estimated for carbon-black-filled EPDM. The <Mad> value obtained for

Figure 10.10 Schematic representation of the physical network structure in a carbonblack-filled elastomer [62]. The symbol … indicates elastomer - carbon black adsorption junctions. The length scales in this figure and the EPDM/carbon black volume ratio are fictional. For simplicity, none of the contacting carbon black aggregates, which form agglomerates, have been included Reprinted with permission from V.M. Litvinov and P.A.M. Steeman, Macromolecules, 1999, 32, 25, 8476, Figure 6. Copyright 1999, American Chemical Society

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Spectroscopy of Rubbers and Rubbery Materials the bound EPDM rubber was about 1800-2500 g/mol and was found to depend on carbon black aggregation and on the content and type of filler [62]. The mean end-toend distance between adsorption junctions was comparable to the average distance between adjacent carbon black aggregates [62]. This suggests that carbon black aggregates in the bound rubber are interconnected by EPDM tie chains, and a continuous EPDM/ carbon black physical network is formed in the bound rubber. It was shown that the physical EPDM-carbon black network is largely responsible for determining the modulus at high strains [62]. The results obtained for unvulcanised EPDM and NR filled with carbon black provide convincing evidence that the physical network has a ‘bimodal’ structure [62, 79]. Two types of EPDM chains and/or chain fragments with widely differing densities of EPDMcarbon black adsorption junctions are present in the rubbery matrix outside the EPDMcarbon black interface (tightly bound rubber) (Figure 10.11) [62]. The first is an EPDM fraction which is loosely bound to the carbon black due to adsorption interactions. This loosely bound rubber has numerous adsorption network junctions, similar to those in bound rubber. The second EPDM fraction, consisting of extractable rubber, contains a relatively small number of adsorption network junctions and can apparently be extracted from the compounds. The fraction of loosely bound EPDM chains determined with the aid of NMR increases with an increase in the maximum possible EPDM-carbon black contact area per unit volume of the elastomer, regardless the type of carbon black used, and is relatively close to the content of bound rubber [62]. Several authors have studied the effect of various structural characteristics of carbon black on the rubber-filler interactions. They found that the interfacial fraction is larger at rough carbon black surfaces with numerous edges and ledges [93]. This suggests that edges and ledges are the active sites whose surface energy is quite a lot higher than that of a smooth quasi graphitic scales of carbon black [94]. The strength of rubber-filler interactions can be varied by modifying the carbon black surface by means of oxidation or chemical treatment. Such treatments result in the formation of hydroxyl and acid groups at the carbon black surface. The mobility of the NR chains at the modified carbon black surface is greater than that at untreated surfaces, which suggests that such modification results in a decrease in the number of active sites and the strength of the adsorption, resulting in the observed increase in mobility [95-97]. Oxidation moreover causes a change in the structure of the physical network in the rubbery matrix. Selective 1H T2 relaxation experiments were used to determine specific sites of styrenebutadiene (SBR) rubbery chains, which interact with the carbon black surface [84, 98, 99]. The T2 of all the different kinds of protons in SBR chains was found to consist of at least two components attributable to a highly immobilised, tightly bound rubber around the filler particles and an outer region of loosely bound and free rubber. The relative

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Figure 10.11 The structure of an EPDM/N550 (phr=100) vulcanisate according to the results of NMR and extraction studies - (A) and mechanical data for the case of pure hydrodynamic reinforcement - (B) [62]. The volume fraction of microphases/ components is given in vol.%. According to the NMR data, the total network density in the rubber phase, 1/2Mc+e+ad, equals 425 mmol/kg, where subscripts c, e and ad stand for chemical crosslinks, chain entanglements and adsorption rubber-filler junctions. The density of the adsorption junctions in the loosely bound rubber, 1/ 2Mad, equals 17 mmol/kg Reprinted with permission from V.M. Litvinov and P.A.M. Steeman, Macromolecules, 1999, 32, 25, 8476, Figure 18. Copyright 1999, American Chemical Society

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Spectroscopy of Rubbers and Rubbery Materials immobilisation of different types of chain units at the carbon black surface suggests that double bonds of butadiene chain units are preferentially adsorbed by carbon black, unlike the aromatic and methylene chain fragments. Advanced 2-dimensional NMR experiments will have to be carried out to determine the conformational state of polymer chains at the carbon black surface and the exact nature of the rubber-filler interactions. Several T2 relaxation studies have shown that the fraction of rubber-filler interface and the structure of the physical network is affected by dispersion of carbon black in the rubbery matrix, which is improved with increasing the length of the mixing time [62, 100-104]. In the first stage of mixing, the fraction of the rubber-filler interface and the density of rubber-filler physical junctions increase considerably with the length of the mixing time. The mobility of rubber chains at the interface decreases rapidly upon mixing [103,104]. In the second stage, the amount of interface increases at a slower rate, but the chain mobility becomes independent of the mixing time. In the third stage, no appreciable change is observed in the amounts of each component. The tightly bound rubber appears to be formed very quickly when the rubber segments come into contact with a new active carbon black surface, resulting in rapid growth of the loosely bound rubber component around the rubber-filler interface.

10.5.3 Silica-Filled Silicon Rubbers Physical interactions between PDMS and silica surfaces have been studied in close detail using model compounds containing a large fraction of silica with a large specific surface [88, 105-113]. Using 1H, 2H, 13C T1 and T2 relaxation experiments, unambiguous information on the short-range (local) chain dynamics near the surface of hydrophilic and silylated silicas was obtained in these studies. 1H and 2H spin-lattices or longitudinal (T1) magnetisation relaxation experiments show two types of PDMS chain units with widely differing local chain mobilities [88, 105, 109]. 1H T1 data obtained for PDMS containing different amounts of a fumed silica are shown in Figure 10.12. The frequency of the chain motions of one of the two types of chain units, inferred from the minimum T1 around 195 K (–78 °C), is similar to that observed in unfilled PDMS. The mobility of the chain units of the other type is greatly hindered, resulting in a T1 minimum at higher temperature (at about 7 °C). This minimum becomes more pronounced in proportion to the total surface of filler particles in the mixture. It was suggested that the observed chain immobilisation is caused mainly by physical adsorption of PDMS chain units adjacent to the silica surface and to a lesser extent by entropy constraints at the silica surface [113]. The estimated thickness of the PDMS-silica interface is about 1 nm [105, 109, 110], which is comparable with that of carbon-black-filled conventional rubbers. The adsorption

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Figure 10.12 Temperature dependence of the 1H T1 relaxation time of unfilled and filled PDMS [113]. Experimental values are indicated by points and the contributions to T1 deriving from the C3- rotation of CH3 groups [T1 minimum at ~ 100 K (-173 °C)], from chain motion in the PDMS matrix and at the PDMS-silica interface (T1 minima at ~ 195 K (-78 °C) and ~ 280 K (7 °C), respectively) are indicated by solid lines. Also indicated is the volume fraction of hydrophilic silica (the specific surface area is 286 m2 g-1) in mixtures. The experiments were performed on PDMS-silica mixtures, which were prepared by adding silica to a dilute PDMS solution in CCl4 followed by solvent evaporation after prolonged stirring of the suspension. The measurements were performed using amorphisised samples [113] Reproduced with permission from V.M. Litvinov in Organosilicon Chemistry II. From Molecules to Materials, Eds., N. Auner and J. Weis, VCH, Weinheim, Germany, 1996, 779, Figure 6. Copyright 1996, VCH Publishers

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Spectroscopy of Rubbers and Rubbery Materials energy per single SiO(CH3)2 bond at the surface of hydrophilic silica is about 8-12 kJ/ mol [108]. It was further suggested that adsorption of polymer chains results in extended chain conformation at the silica surface [113, 114]. The total energy per binding site, which involves 10-15 hydrogen bonds between PDMS chain units and the surface of hydrophilic silica, was estimated to be about 120 kJ/mol [115, 116]. 13

C NMR T1 and T2 relaxation experiments using adsorbed PDMS at submolecular coverage have shown that chain dynamics within the interface is dependent on the extent of surface coverage [110, 111]. It was shown that the chain mobility increases as the chain surface density increases. It was suggested that these changes reflect the conformational state of the adsorbed chains. At a very small degree of surface coverage the chains apparently adopt an extremely flat profile, with few loops or tails longer than approximately 0.5-1.5 nm. As the surface density (degree of surface coverage) increases above two monolayers, the polymer chain segments are driven from the surface into loop and tail configurations. The degree of immobilisation at a fixed distance from the surface is determined by the size of the loops and tails. The mobility of loops and tails increases with increasing their length, because of a decrease in the number of constraints. It was suggested that the backbone architecture is the key factor in restricting local chain mobility at a surface. Polymer chains of greater stiffness would probably show both a larger boundary region and a greater decrease in chain mobility at the filler surface.

The relative number of immobilised adsorbed chain units in a single type of silica is proportional to Mn-0.5 [117]. The fraction of immobilised chain units is larger in hydroxylterminated PDMS [117]. This suggests that the majority of hydroxyl chain ends are linked to the silica surface due to the formation of double hydrogen bonds. Contrary to carbon-black-filled conventional rubbers, which form a semi-rigid interface at the carbon black surface, PDMS chain units at the silica surface are not rigidly linked to the silica surface. Two types of dynamic processes are thought to occur at the interface: relatively fast anisotropic reorientation of chain units in the interfacial layer and slow adsorption-desorption of chain units (Figure 10.13) [108, 113]. When the temperature increases, both the fraction of less mobile, adsorbed chain units and the lifetime of the chain units in the adsorbed state decrease. The lifetime of PDMS chain units in the adsorbed state approaches zero at approximately –73 °C and 227 °C at the surface of hydrophobic and silylated silicas, respectively [108, 113]. The frequency of PDMS chain mobility in the interfacial layer is largely dependent on the type of silica surface. Water adsorption by the hydrophilic silicas causes the degree of chain immobilisation in the interface to decrease [107]. Sililation of the silica surface results in a significant decrease in the strength of adsorption interactions at the silica

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Figure 10.13 Schematic representation of an adsorbed fragment of a PDMS chain [113]. The solid arrows denote the anisotropic motion of chain units during the lifetime in the adsorbed state. The open arrows denote the adsorption-desorption of chain units. The length of the adsorption-desorption jumps is possibly in the order of 0.1-0.5 nm Reproduced with permission from V.M. Litvinov in Organosilicon Chemistry II. From Molecules to Materials, Eds., N. Auner and J. Weis, VCH, Weinheim, Germany, 1996, 779, Figure 8. Copyright 1996, VCH Publishers

surface, which causes a decrease in the fraction of the interface and a significant increase in chain dynamics in the interfacial layer [107, 109, 113]. A good understanding of the structure of the network in filled rubbers is of great importance, because the rubber’s elastic properties are determined primarily by the density of chemical and physical network junctions and their ability to fluctuate. The following types of network junctions occur in filled rubbers: 1) chemical crosslinks between rubbery chains and between rubbery chains and the filler surface, 2) adsorption junctions at the filler surface, 3) temporary and trapped chain entanglements, and 4) topological hindrances of elastomeric chains due to the confinement of chains in a restricted geometry or elastomer-filler entanglements [113, 118]. The relative contribution of each of these types of network junctions to the overall crosslink density in silica-filled PDMS was estimated by means of 1H T2 relaxation experiments [113, 118-121]. To determine the relative contributions of the different types of network

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Spectroscopy of Rubbers and Rubbery Materials junctions to the overall network density, the relaxation experiments were performed using original samples, swollen samples and samples swollen in the presence of ammonia. In the case of the original samples all the different types of network junctions were found to contribute to the overall network density. The effect of temporary chain entanglements and topological constraints diminishes largely in the swollen samples. Since ammonia is preferentially adsorbed onto the silica surface [113] its molecules cause desorption of physically adsorbed PDMS chains. In this case the network density was determined mainly by chemical crosslinks and trapped chain entanglements. Combined analysis of the T2 data obtained for the same sample under three different conditions showed that adsorption junctions and topological constraints are largely responsible for the overall crosslink density in the case of silica-bound PDMS. The apparent number of elementary chain units between the topological hindrances in bound rubber is estimated to be approximately 40-80. As in carbon-black-filled EPDM and NR rubbers, the physical network in silica-filled PDMS has a ‘bimodal’ structure [61]. A loosely bound PDMS fraction has a high density of adsorption junctions and topological constraints. Extractable or free rubber does virtually not interact with the silica particles. It was found that the density of adsorption junctions and the strength of the adsorption interaction, which depends largely on the temperature and the type of silica surface, largely determine the modulus of elasticity and ultimate stress-strain properties of filled silicon rubbers [113].

10.5.4 Silica-Filled Conventional Rubbers The use of silicas for reinforcing conventional rubbers is receiving ever more attention due to the specific properties of the composites, in particular for their use in tyres [122, 123]. Hydrogen-bond interactions between surface silanol groups in silica agglomerates are very strong, in particular in comparison with the interactions that take place between the silanol groups of the silica and the commonly used non-polar olefinic hydrocarbon rubbers. This makes it very difficult to mix silica with rubber. Chemical modification of the silica surface using functional organosilanes enhances the compatibility of hydrocarbon rubbers and silica, leading to marked improvements in the mechanical properties of silica-reinforced rubber vulcanisates. 1

H NMR transverse magnetisation relaxation experiments have been used to characterise the interactions between NR, isoprene rubber, BR, EPDM and polyethylacrylate rubbers with hydrophilic silica and silicas modified with coupling agents [124-129]. These studies showed that the physical interactions and the structures of the physical networks in rubbers filled with carbon black and rubbers filled with silicas are very similar. In both cases the principal mechanism behind the formation of the bound rubber is physical adsorption of rubber molecules onto the filler surface.

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Characterisation of Chemical and Physical Networks in Rubbery Materials Using … The mobility in both tightly and loosely bound BR and isoprene rubbers increases, and the fraction of bound rubber decreases with a decreasing concentration of silanol groups on the silica surface [124]. This led to the suggestion that the silanol groups on the silica surface are active sites for the chain adsorption. The grafting of aliphatic chains to the silica surface leads to a decrease in BR-silica interactions [125]. The effect is less pronounced in BR filled with carbon black containing aliphatic chains at the surface. The number of immobilised rubbery chains on the silica surface seems to be dependent on the coupling agent used and the grafting density. A significant number of immobilised NR chains were found on the surface of hydrophilic silica and silica grafted with 3mercaptopropyl triethoxysilane (MPTES) [128, 129]. Strong NR-filler interactions were observed at 100 °C even in the presence of a good solvent, suggesting that a quasipermanent rubber-silica network formed in the samples due to physical adsorption of chain units at the silica surface. Mixtures of NR filled with silicas that were treated with bis(triethoxysilyl propyl)tetrasulfide (TESPT) and triethoxysilyl benzothiazole (TESBT) contain only a small amount of immobilised NR [128, 129]. Since the fraction of immobilised chains is determined by the strength of the adsorption interactions and the degree of dispersion of silica particles in the rubber matrix, NMR experiments using well-defined model compounds will have to be carried out to obtain a good understanding of the effect of silica treatment on rubber-filler interactions.

10.6 Chains Grafted onto a Filler Surface Chain grafting is widely used to modify the surface of active fillers and to improve filler dispersion in the matrix material. Fumed silicas are used in the reinforcement of rubbers and the thickening of polymeric liquids. By varying the grafting density, its heterogeneity and the average length of the grafted chains, the silica-silica interactions and interactions between silica particles and the matrix material can be influenced to a large extent. Therefore, knowledge of the molecular structure of the grafted layer is necessary for a proper understanding of the molecular mechanisms responsible for the reinforcing and thickening effects. The molecular mobility of high-molecular-mass PDMS chains grafted to the silica surface at one end has been studied with the aid of various solid-state NMR techniques [130-132]. The use of 2H NMR methods for this purpose is reviewed in the present book [133]. A PDMS layer chemically attached at both chain ends to the surface of hydrophilic silica has been studied by means of 1H T2 relaxation experiments [112]. The average length of the grafted chains in samples studied was varied from about 4 to 8 Si-O bonds. These experiments showed that the T2 relaxation method is very sensitive to the heterogeneous

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Spectroscopy of Rubbers and Rubbery Materials mobility of PDMS chains on a silica surface. Since molecular mobility is closely coupled to chain length, grafting density and chain length distribution, T2 experiments can provide valuable information on the structure of grafted layers. The grafted PDMS layer was found to consist of immobilised chain segments at the PDMS-silica interface and mobile chain portions outside the interface similar to that of adsorbed PDMS. The interface fraction increases proportionally with a decreasing chain length (Figure 10.14). About four dimethylsiloxane pendent chain units next to the grafting site were immobilised due to chain anchoring to the silica surface.

Figure 10.14 The fraction of the PDMS-silica interface, as determined by the fraction of the T2 relaxation component (%T2in) with a decay time of about 80 μs, against the length of the grafted chains, Nt, in silica samples swollen in C2Cl4 in the presence of NH3 [112]. The Nt value corresponds to the total length of any chains grafted at one end and the halflength of chain loops and chains attached to neighbouring silica particles. The line represents the result of a linear regression analysis of the data using the following equation: %T2in = 1 – (Nt – Nad)/Nt, where Nad is the number of immobilised -SIO(CH3)2chain units at the PDMS-silica interface. The correlation coefficient equals 0.89 Reproduced with permission from V.M. Litvinov, H. Barthel and J. Weis, Macromolecules, 2002, 35. Copyright 2002, American Chemical Society

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Characterisation of Chemical and Physical Networks in Rubbery Materials Using … The fraction of immobilised PDMS chain units decreased slightly when the swollen samples were treated with ammonia, which molecules cleft adsorption bonds between PDMS chain units and silanol groups at the silica surface [113]. This suggests that a small fraction of -SiO(CH3)2- chain units was immobilised as a result of physical adsorption at the silica surface. The fraction of physically adsorbed chain units appeared to be proportional to the number of residual silanol groups on the silica surface. The mobility of the chain portions outside the interface was found to differ significantly in the different samples studied, and to increase with an increasing average length of the grafted chains. The NMR method allowed us to distinguish a grafted layer with a dense ‘brush-like’ structure containing chains of a fairly uniform length and a layer containing a significant fraction of long chain loops.

10.7 Semi-Crystalline Elastomers Rubbery chains in various rubbery materials, such as elastomeric polyolefins and EPDM rubbers with a high ethylene content, form a crystalline micro-phase. Chain anchoring to crystals leads to the formation of a physical network, with the crystallites acting as multifunctional crosslink junctions [134-136]. Above the Tg, the mean chain length between anchoring points to the crystalline phase largely determines the T2 relaxation time of the amorphous phase [136-139]. The effect of chain confinement of amorphous segments by crystallites moreover induces additional constraints of segmental motions in the amorphous phase [139]. At a low degree of crystallinity, T2 relaxation experiments can provide information on the mean molecular mass of polymer chains interconnecting adjacent crystals [136, 138]. Wide-line 1H and 2H NMR spectra and T2 relaxation experiments have been used to determine the composition of the phases in semi-crystalline polymers [133, 136, 138144]. The experiments were also used to obtain real-time information on the kinetics of crystallisation and melting [143-148]. The use of high-resolution NMR methods to characterise semi-crystalline polymers is reviewed elsewhere [17, 18, 30, 34, 149]. The T2 data obtained for semi-crystalline polymers are often analysed with the aid of a three-phase model comprising a crystalline phase, a semi-rigid interphase2 and a soft

2

The degree of molecular mobility in the interphase lies between that in the rigid crystalline phase and that in soft amorphous phase. The fraction of the intermediate phase was found to depends on the employed technique, the temperature and the method of data evaluation. In many respects, the intermediate phase has a kinetic origin and may not be regarded as a true thermodynamical phase. It would apparently be more correct to define the third phase as an interface or a semi-rigid fraction of the amorphous phase.

381

Spectroscopy of Rubbers and Rubbery Materials amorphous phase. Good correlation was observed between crystallinity measured by means of NMR and with the aid of other methods [140-144]. The fraction of highly disordered nano-size crystals can also be determined with the aid of the T2-relaxation method [138]. In the case of elastomeric materials with a low degree of crystallinity a simple two-phase model can yield a good description of the T2 relaxation behaviour, as can be seen in Figure 10.15 [136, 138]. This figure shows that the 1H T2 decay for ethylene/1-octene copolymer can be well described with two components, whose characteristic decay time is typical for crystalline and amorphous phases. It appears in this case that the two-phase model may yield reliable results, because 1) the mobility of chains adjacent to the surface of small crystals is greater than that at the surface of large crystals due to the larger disorder at the surface of small crystals and crystal disorder that promotes chain mobility, and

Figure 10.15 The decay of the transverse magnetisation (points) for ethylene-octene copolymer at different temperatures [136]. The decay was measured using the solidecho pulse sequence. The solid lines represent the result of a least-squares adjustment of the decay using a linear combination of Weibull and exponential functions. The dotted lines represent the relaxation component with a long decay time. In the experiments the sample was heated from room temperature to 343 K (70 °C) Reproduced with permission from V.M. Litvinov, Macromolecules, 2001, 34, 24, 8468. Copyright 2001, American Chemical Society

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Characterisation of Chemical and Physical Networks in Rubbery Materials Using … 2) the broad distribution of the correlation times of chain motions in the amorphous phase can cause ‘apparent’ single exponential relaxation.

10.8 Ionic Viscoelastic Materials Polymers with small concentrations of ion-containing groups (up to about 15 mol%) are known as ionomers, as opposed to polymer electrolytes, which contain larger amounts of ionic groups. Elastomers containing ionic groups can be used as thermoplastic rubbers. The presence of ionic groups in polymer chains greatly affects their physical properties [150-152]. Ionic aggregation causes a marked increase in viscosity and the strength of such polymer materials. The main mechanism responsible for the increase in strength is dependent on strong electrostatic interactions between ionic groups, resulting in the formation of immobilised ionic multiplets and their aggregates and clusters. Polymer chains interconnect multiplets, resulting in the formation of a physical network (Figure 10.16) [153].

Figure 10.16 A schematic representation of the morphology of a telechelic ionomer where (1) represents the ion-containing core, (2) an immobilised interfacial layer and (3) a phase containing network chains, dangling chains and free chains [153] Reprinted with permission from V.M. Litvinov, A.W.M. Braam and A.F.M.J. van der Ploeg, Macromolecules, 2001, 34, 3, 489, Figure 17. Copyright 2001 American Chemical Society

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Spectroscopy of Rubbers and Rubbery Materials Multiplets are supposed to contain a relatively small number of ionic pairs and consist of ionic material only, whereas clusters are large structures that are rich in ion pairs and contain a considerable quantity of an organic material. Recent proton NMR relaxation studies have shown that the method can be used to study the formation of ionic clusters and the structure of the physical network in ionomers [153].

10.9 Rubbery Phases in Blends and Emulsions The most comprehensive information on the phase structure and molecular mobility in polymer materials can be gained with high-resolution NMR techniques [17, 18, 30, 31, 34]. Low-resolution NMR techniques may offer some advantages in the following cases: 1) characterisation of phase structure and network density in polymers with a complex chemical composition that result in substantial overlapping of resonances from different chemical groups, hampering the interpretation of high-resolution NMR data, 2) fast quantitative analysis of phase composition, phase transition and network structure, and 3) when accurate control of sample temperature is required, which can be easily realised with the aid of low-field NMR equipment, allowing analysis of phase transitions. Proton NMR magnetisation relaxation experiments are often used to determine microphase structure and molecular motions in polymer materials containing rubbery phases [17]. The results of proton NMR transverse magnetisation relaxation experiments for heterogeneous materials are easier to interpret than those of T1 and T1ρ (spin-lattice relaxation in the rotating frame) experiments, because 1H T2 data are not complicated by the effect of spin-diffusion [17, 18]. Moreover, the T2 relaxation process is very sensitive to even small changes in molecular mobility, as can be inferred from the T2 range from about 10 μs observed for glassy and crystalline phases/components to a few seconds measured for low-molecular-mass liquids. This great difference in T2 shows that this experiment is highly sensitive to changes in the chemical structure, composition and morphology of polymer materials. However, care should be taken in analysing T2 relaxation decay curves with respect to the different phases/components contained in the polymer (see Section 10.3). Analysis of 1H T2 relaxation as a function of temperature yields information on molecular motions. Side-groups and local chain motions (secondary relaxations) cause a change in T2 below the Tg. The value of T2 increases with the amplitude and the frequency of molecular motions. The glass transition that occurs in the time scale in the NMR T2 relaxation

384

Characterisation of Chemical and Physical Networks in Rubbery Materials Using … experiments, i.e., 10-100 μs, causes a sharp increase in T2. The Tg measured in such NMR experiments is comparable with the dynamic Tg measured by means of DMA at frequencies of about 10-50 kHz and is usually observed at a temperature 30-50 °C above Tg measured with differential scanning calorimetry. T2 experiments are used to determine the concentration of (inter)phases/components in polymer materials if polymer chains in these (inter)phases/components reveal a significant difference in molecular mobility [17, 34]. In such cases, the T2 relaxation function is the weighted sum of the T2 decays of different components/phases. The relative fraction of these components is proportional to the concentration of hydrogen in these (inter)phases/ components (see Sections 10.4 and 10.7). The characteristic decay time, T2, is related to molecular mobility in different phases. Several examples of such studies are given in this section. The 1H T2 relaxation method has been used to determine the crosslink density in the rubbery phases in (styrene)-b-(butadiene)b-(styrene) tri-block copolymers (SBS) [154, 155], BR rubber in polystyrene (PS) used to improve impact strength [156] and acrylic rubber in impact-modified a styrene acrylonitryl copolymer (SAN) [157]. It was found that a high degree of toughness can be obtained only if the rubber phase has a low crosslink density, resulting in low cavitation resistance of the rubber particles under strain [157]. 1

H T2 experiments have also been used to see whether they are suitable for analysing interfaces between different phases in polymer blends, such as SBS block copolymers [158] and (methylmethacrylate)-b-(n-butylacrylate) block copolymers [159]. The distribution of crosslinks in blends of different rubbers has been determined by means of line-width analysis of 1H high-resolution NMR spectra obtained using swollen samples [33,160-166]. It should be noted that the results of these experiments suffer from systematic errors, since the line width of the NMR spectra is affected not only by the crosslink density, but also by differences in the magnetic susceptibility of the different phases [79-83] and even by the effects of slow sample rotation [20].

1

H T2 relaxation experiments have moreover been used to investigate the phase composition, interfaces and phase diagrams of polymer solutions, emulsions and coreshell lattices [167-173].

10.10 Real-Time NMR Experiments Two practical aspects are of importance in vulcanising rubber compounds: 1) the crosslinking kinetics, which determine the production cycle time or throughput, and

385

Spectroscopy of Rubbers and Rubbery Materials 2) the crosslink density, which partly determines the mechanical and elastic properties of the vulcanisate. The crosslinking kinetics and the final state of cure are commonly studied with the aid of rheometers. NMR relaxation experiments can offer several advantages for the characterisation of the crosslinking kinetics in complex materials because of high method selectivity with respect to the rubbery chains/phases in polymer blends, filled and oil extended rubbers. Network structures have been quantitatively determined by means of real-time 1H NMR T2 relaxation experiments for several polymers [174-178]. The effect of the curing conditions on sol and gel fractions and the spatial heterogeneity of the network structure has been studied for polyethylene [174], polyacrylamide [175], PDMS [176], BR [177], epoxy resins [178] and EPDM [179]. A change in 1H T2 relaxation during accelerated sulfur vulcanisation of EPDM rubber is shown as an example in Figure 10.17 [179].

Figure 10.17 Dependency of T2 relaxation time as a function of the actual vulcanisation time during accelerated sulfur vulcanisation of EPDM [179]. The line has been included to guide the eye. The vulcanisation temperature was 413 K (140 °C) Reproduced with permission from V. Litvinov and M. van Duin, Kautschuk und Gummi Kunststoffe, 2002, 55. Copyright 2002, Huthig Verlag

386

Characterisation of Chemical and Physical Networks in Rubbery Materials Using … The time it takes for the temperature of a sample to stabilise in an NMR probe can also be determined in real-time 1H NMR T1 relaxation experiments, because at vulcanisation temperatures T1 is only slightly affected by a moderate crosslink density [22, 180, 181]. The T2 data obtained in the aforementioned experiments were used for determining an increase in the density of chemical crosslinks upon vulcanisation time (Figure 10.18) [179]. The method can also be used in kinetic studies of the vulcanisation of filled and oil-extended rubbers.

Figure 10.18 Dependency of the chemical crosslink density (1/2Mc) on the actual vulcanisation time during accelerated sulfur vulcanisation of EPDM [179]. The line has been included to guide the eye. The vulcanisation temperature was 413 K (140 °C) Reproduced with permission from V. Litvinov and M. van Duin, Kautschuk und Gummi Kunststoffe, 2002, 55. Copyright 2002, Huthig Verlag

10.11 Low-Field NMR Magnetisation Relaxation Experiments for Quality Control Purposes The high selectivity of solid-state NMR with respect to different phases/components in polymer materials can offer several advantages for quality tests over other analytical techniques. Although access to information on the use of NMR for quality-control purposes is rather limited, several applications for fast analysis of rubbery materials have been published, for example for determining the concentration of oil in rubbers

387

Spectroscopy of Rubbers and Rubbery Materials [182, 183], the distribution of carbon black in rubber matrices [100,101], the solids content of rubber lattices [184], the viscosity of oil [185], the crystallinity of polyolefins [186], molecular masses and melt flow indices [186] and the component contents of copolymers and polymer blends [186]. The method is also used for in-process analysis in polymer industry [186, 187]. The NMR method was found to offer the following advantages for quality tests: • It yields quantitative data on network density and phase composition • It shows a high selectivity with respect to phases/components in blends • Mineral fillers do not interfere with the results • Samples do not have to meet any special requirements; swollen gels, solid materials, emulsions, suspensions, granules and powders can all be analysed • The method measures volume-average properties • It yields results quickly • It is environmental friendly because small amounts of chemicals are used. Quality tests are usually performed using bench-top, low filed NMR spectrometers. Volume-average properties are determined with this equipment. Surface layer of samples can be analysed using recently developed NMR-MOUSE (mobile universal surface explorer) [26, 188]. The NMR-MOUSE is a relatively small NMR device suited for the investigation of surface-near volume elements. Lateral surface heterogeneity of elastomeric materials can be scanned with this device. Possible applications of the NMR-MOUSE for the characterisation of rubbery materials were demonstrated [26,189-191].

10.12 Conclusions We may safely conclude that proton NMR magnetisation relaxation experiments are a valuable tool for characterising various types of viscoelastic materials and rubbery phases in blends, emulsions and core-shell lattices. The NMR method can be used to quantitatively analyse the structure of chemical and physical networks as well as network defects. The great advantage of the NMR method over traditional macroscopic methods is due to the high selectivity of the NMR experiments with respect to different (inter)phases/components in materials intended for practical use. This information is essential for obtaining a better understanding of structureproperty relationships. NMR relaxation experiments may also be more advantageous in routine quality control than traditional methods.

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11

High Resolution NMR of Elastomers Yasuyuki Tanaka, Eng Aik-Hwee and Jitladda Sakdapipanich

11.1 Introduction The structure of elastomers gives an important information to polymer scientist and technologist because it is closely related to the properties and hence performance of the polymer. The information on the structure of elastomers also provides important clues on the mechanism of polymer synthesis, or biosynthesis in the case of naturally occurring elastomer. This knowledge would certainly contribute to a better control in producing elastomers with certain desirable properties. Therefore, structural characterisation of elastomers, whether natural or synthetic, is a subject of paramount important in many research works. In this aspect, high-resolution nuclear magnetic resonance (NMR) is one of the most powerful tool for the structural characterisation, because of its high sensitivity, applicability for quantitative measurements without calibration curve, and versatility applications to elastomers in solid, solution, latex and swollen states. The structure of elastomers can be very complex but it must satisfy two basic requirements, i.e., low glass transition temperature and vulcanisability. In the case of synthetic elastomers, the former can be fulfilled by copolymerisation of common monomers to give higher molecular mobility. There are many diene homopolymers that can satisfy these requirements. Diene copolymers contain double bonds that serve as a crosslinking point for vulcanisation with sulfur or peroxides. The sequence distribution of comonomer units in these copolymers has to be controlled to provide desirable properties as elastomers. With increasing demand in the versatility of physical properties, the structure of copolymers has now been made more and more complicated. For example, styrenebutadiene copolymer (SBR) synthesised by emulsion polymerisation has been recognised to be a random copolymer having good properties for general purpose. However, recent progress in anionic polymerisation makes it possible to produce SBR of a controlled sequence distribution having the random, partial block, or S-B-S triblock structures and the corresponding partial hydrogenated products. These copolymers exhibit excellent properties satisfying individual requirements. More recently, it has been substantiated that certain functional terminal groups of SBR could improve the physical properties of the copolymer dramatically.

401

Spectroscopy of Rubbers and Rubbery Materials In the case of naturally occurring elastomers, such as natural rubber, it is necessary to get a thorough understanding of the biosynthesis process in the plants. The process could be eventually controlled through genetic engineering so that a highly desirable elastomer could be obtained. Knowledge on structure of these elastomers that have certain superior properties also enables scientists to produce the synthetic analogue with the similar properties. High-resolution NMR has been applied to the characterisation of these elastomers in many areas, such as chemical composition analysis, extent of chemical modification, isomeric structure of diene monomer units, sequence distribution, small amounts of modified structure, end-group, etc. However, this technique is less useful in the analysis of cured rubber samples, long block sequences, inhomogeneity of chains such as chemical composition distribution, and very small amounts of structural irregularity. Recent progress in solid-state high-resolution NMR has partly solved the first problem. Highresolution NMR at 500-750 MHz has been shown to improve the sensitivity dramatically and has enabled determination of very small signal derived from terminal group or irregular structures. Valuable reviews and books of structural analysis of elastomers have been published by several authors [1-6]. Some of these reviews provide excellent explanation on the basic theory of sequence distribution of copolymer and NMR techniques applicable to elastomers. Typical high-resolution 1H- and 13C-NMR spectra of various vulcanisates and raw rubbers are depicted in a book written by Kelm [6]. The assignments and references shown for each rubber are very useful for structural studies of elastomers. In view of recent progress in the hardware and software of NMR, this chapter describes some of the more recent applications of high-resolution NMR to the structural characterisation of elastomers, after a brief description on the fundamental structural features of elastomers.

11.2 Structural Feature of Elastomers The analysis of chemical composition of copolymers can be done without difficulty using 1 H-NMR and 13C-NMR techniques, provided that the signals of each component are assigned correctly. However, most of the signals in copolymers reflect the sequence distribution of comonomer units from diad to tetrad or longer sequences. NMR instrument with magnetic field above 500 MHz for 1H-NMR and 125 MHz for 13C-NMR has facilitated the detailed assignment of the signals. In this case the chemical composition is obtained by using the sequence equations such as diad and triad. Diene monomers such as butadiene, isoprene, 2-chlorobutadiene give rise to the following isomeric structures in the polymers shown in Figure 11.1. It should be noted that the

402

High Resolution NMR of Elastomers

Figure 11.1 Microstructures of polybutadiene and polyisoprene

configurational sequences of 1,2 and 3,4 units reflect the arrangement of R and S configuration such as RR or SS (meso) and RS or SR (racemic or racemo) configurations. The ratio of these isomeric structures, which is sometimes referred to as microstructure, can be determined by 1H-NMR and 13C-NMR. For example, the microstructure of polybutadiene can be measured by 1H-NMR with magnetic field higher than 300 MHz or 13C-NMR with magnetic field higher than 25 MHz. The microstructure of polyisoprene can be analysed using similar techniques. However, most of synthetic diene polymers are a kind of copolymers composed of these isomeric units. Signal-split reflects the sequence distribution of three to four isomeric units, configurational sequence of 1,2 and/or 3,4 units, and head-to-tail arrangements. Rubber from Hevea brasiliensis is a typical naturally occurring cis-1,4 polyisoprene termed as natural rubber. Many higher plants and fungi produce cis-1,4 polyisoprene. Natural polyisoprenes appear to be a simple polymeric system, considering that they are produced by enzymatically controlled polymerisation process. However, the reason why only natural rubber exhibits outstanding properties has yet to be clarified. The analysis of detailed structure of natural rubber, in comparison with other natural cis-1,4 polyisoprene, could provide direct information on the origin of the characteristic properties of natural rubber. Gutta percha and Balata from Paraquim gutta and Mamusops balata, respectively, are typical naturally occurring trans-1,4 polyisoprene. Chicle resin from Achras zapota contains both cis-1,4 and trans-1,4 polyisoprenes. Biochemical studies since 1950s have not elucidated the biosynthesis mechanism on the formation of cis-1,4 and trans-1,4 polyisoprenes. The detailed structural analysis would disclose the longstanding puzzles on the biosynthesis mechanism of these natural polyisoprenes.

403

Spectroscopy of Rubbers and Rubbery Materials The analysis of end-groups and irregular structures, e.g., branch-points and grafting points, is a difficult task because of the high molecular weight of elastomers. Elastomers usually have higher molecular weight than ordinary plastics except for some thermoplastic elastomers, such as S-B-S triblock copolymers of styrene and butadiene. For such analysis, a high magnetic field NMR is required in order to obtain a high signal-to-noise ratio. In addition to the sequence distribution, copolymers usually have the distribution of chemical composition, i.e., the difference of chemical composition among polymer chains. Unfortunately, conventional NMR techniques are not sensitive to the chemical composition distribution, because NMR signals reflect an average structure of each polymer chain. Therefore, solvent cross-fractionation or HPLC separation technique followed by composition analysis of each fraction is usually adopted for this purpose. However, the development of on-line high performance liquid chromatography (HPLC), including size exclusion chromatography (SEC), makes it possible to analyse the difference of structural features among polymer chains. In the following section, certain topics are selected to introduce the recent applications of high-resolution NMR to the structural characterisation of elastomers, together with basic information on signal assignments.

11.3 Analysis of Chemical Composition and Sequence Distribution

11.3.1 Accuracy of NMR Measurements of Chemical Composition A series of works has been made on the reliability of NMR data; chemical shifts, intensity and resolution [7], spin-lattice relaxation time and nuclear Overhauser enhancement factor (NOE) [8,9], and composition and comonomer sequence distribution [10], by round robin method using poly(methyl methacrylate) (PMMA), solanesol, and copolymer of MMA and acrylonitrile (AN). The standard deviations (σ) of chemical shift measurements for 1H- and 13C-NMR signals of PMMA were 0.0036-0.0071 and 0.0540.307 ppm, respectively. Here, the measurements were made with 90 to 500 MHz instruments for 1H-NMR. The relative intensities of 3(CH2 + α-CH3)/OCH3 were 4.94 (σ = 0.156), showing a good agreement with the theoretical value of 5. In the case of 13 C-NMR, the relative signal intensities corrected by the NOE value agreed with the theoretical value within ± 10% relative errors, with an appropriate pulse repetition time. The composition of MMA-AN copolymer was determined with the σ value of 3.7-9.5%

404

High Resolution NMR of Elastomers for both 1H- and 13C-NMR measurements. The monomer reactivity ratio for a penultimate model was determined with the σ value 5-14%. The quantitative analysis of end-groups provides information on the number average molecular weight, number of branch-points per chain, and mechanisms of initiation, termination and chain transfer reactions. These are possible only if the structure of end-groups is identified. Direct measurements of the relative intensity for small signals depend on the signal-to-noise ratio and dynamic range of digitiser (ADC). In a model experiment, the determination of CH3NO2 in the solution of PMMA is possible even at a level as low as 1/250,000, with accuracy of 2.5-10.2%, using 1H-NMR at 500 MHz. However, it is necessary to do accumulation of 18,000 scans [11,12]. This will be the limit of quantitative measurement from relative intensity between large and small signals, by considering the dynamic range of ADC. The limit of quantitative measurement can be improved to 1/106 by the use of a coaxial tube, containing a standard sample in the internal tube [12]. Various types of end-groups were analysed quantitatively for PMMA by using 1H-NMR at 750 MHz [13]. One of the advantages of using NMR technique is its ready application to quantitative analysis. Usually, it does not require authentic sample nor calibration curve. However, it is necessary to check the accuracy and precision of the measurement, as mentioned above. It is also useful to check these by comparison with another direct method such as titration method that is applied frequently for the determination of functionality of telechelic diene polymers, as mentioned in the following section. The unsaturation in butyl rubber has been analysed by the epoxy titration method using 3-chloroepoxybenzoic acid or simply iodine index test. The disappearance of all the olefinic signals in butyl rubber was clearly indicated after the epoxy titration, as shown in Figure 11.2. The epoxy titration method showed little deviation such as σ =0.0128 for butyl rubber containing 1.79% unsaturation. However, it is remarkable that a side reaction may occur after reaching the end-point of titration. The unsaturation values with 1H-NMR at 500 MHz showed a good agreement with those obtained by the epoxy titration method as well as those estimated by copolymerisation kinetics [14]. 1H-NMR measurements showed a similar good agreement with analysis by epoxy titration method, by mass balance of monomers, and by predicted method from copolymerisation equations [15]. The molecular composition of ethylene-propylene copolymer (EPM) was analysed using the round robin analysis on a set of 10 EPM samples by 13C-NMR measurements mainly at 75 MHz [16]. According to the signal assignments given in Table 11.1, it is possible to set up several methods to determine the content of E and P. Here, the following four methods were applied [16]. The consistent and reproducible molecular composition results were obtained independent of the measuring laboratory and the following calculation method.

405


Figure 11.2 1H-NMR spectra of butyl rubber, (a) olefinic region, (b) full spectrum after epoxidation, observed at 500MHz [14] Reproduced with permssion from J.E. Puskas and C. Wilds, Rubber Chemistry and Technology, 1994, 67, 329, Figure 1. Copyright 1994, Rubber Division, ACS

Method 1: S = all secondary carbons T = all tertiary carbons P = all primary carbons Mol % ethylene = [3-(S + T - 2 x P)/(S + T)] x100 Mol % propylene = 100 - mol % ethylene Method 2: P′ = Sαα+ 0.5 x (Sαβ + Sαδ ) E′ = 0.5 x [3-Sββ + Sβγ + Sβδ + Sγγ + Sγδ + Sδδ + 0.5 x (Sαβ + Sαγ + Sαδ )] Mol % ethylene = [3-P′/(E′ + P′)] x 100 Mol % propylene = [3-E′/(E′ + P′)] x 100 Method 3:

406


Table 11.1 13C-NMR signal assignment for EPM in tretrachloroethane-d2 [16] Peak No.

Carbon Type

Chemical Shift

Sequence

Integration Limit

1

Sαα

48.1 ÷ 45.3

Sαα

48.50-44.50

2

Sαγ

38.8

r-Sαγ

3

Sαδ

38.4

r-Sαδ

4

Sαγ

37.96

m-Sαγ + m-Sαγ

5

S αδ

37.58

m-Sαδ + m-Sαδ

6

Sαβ

35.7

r-Sαβ

7

Sαβ

34.9

Sαβ + Sαβ

8

Tγγ

33.9

Tγγ

9

Tγδ

33.6

Tγδ

10

T δδ

33.3

Tδδ

11

Tβγ

31.2

Tβγ (m)+ Tβγ(r)

12

T βδ

30.9

Tβδ (m)

13

Sγγ

30.8

Sγγ +Tβδ (r)

31.91-30.61

14

Sγδ

30.4

Sγδ

30.61-30.23

15

S δδ

30.0

Sδδ

30.23-29.32

16

Tββ

28.8 ÷ 28.5

Tββ (mm)+Tββ (mr+rr)

29.15-28.22

17

Sβγ

27.85

Sβγ

28.22-27.63

18

S βδ

27.45 ÷ 26.3

Sβδ

27.63-26.63

19

Sββ

24.9, (24.7, 24.6)

Sββ

25.60-23.95

20

Pββ

22.0 ÷ 21.3

Pββ (mm)

21

Pβγ

21.3 ÷ 20.6

Pββ (mr)+Pβγ(m)+ Pβδ(m)

22

Pγγ

20.6 ÷ 19.5

Pββ(rr)+Pβγ(r)+Pβδ(r)+ Pγγ

40.00-36.50

36.20-34.30

34.29-32.80

22.50-19.00

Reproduced with permission from S.D. Martino and M. Kelchtermans Journal of Applied Science, 1995, 56, 1781, Table 2. Copyright 1995, John Wiley & Sons

407

Spectroscopy of Rubbers and Rubbery Materials N = total no. methylenes = (Sαα+ Sαβ + 3 x Sββ + 2 x Sβγ + 5 x Sγγ + 3 x Sγδ + Sδδ ) N1 = total no. methyls = (Pββ + Pβγ + Pγγ ) Mol % ethylene = [3-(N0 - N1)/(N0 + N1)] x 100 Mol % propylene = [3-2 x N1/(N0 + N1)] x 100 Method 4: n0 = number-average sequence length of uninterrupted methylenes (Sαα+ Sαβ + 3 x Sββ + 2 x Sβγ + 5 x Sγγ+ 3 x Sγδ + Sδδ ) = ⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯ (Sαα+ 0.5 x Sαβ + Sββ + 0.5 x Sβγ + Sγγ+ 0.5 x Sγδ) Mol % ethylene = [3-2/(1 + n0)] x 100 Mol % propylene = 100 - mol % ethylene The composition analysis of SBR, i.e., styrene, cis-1,4, trans-1,4 and 1,2 (vinyl) units, has been carried out by 13C-NMR based on signal assignments indicated in Table 11.2. The integrated intensities of regions A to I are used to calculate percent composition by using following equations [17]. %S = 100 x (A + D)/(A + B + C + D + E) %V = 100 x (B + E)(A + B + C + D + E)

(11.1)

%T = (100 - %S - %V) x (F + G)/(F + G + H + I) %C = (100 - %S - %V) x (H + I)/(F + G + H + I) Determination of the composition of many samples is time-consuming, since integration technique often requires careful check on each spectrum, due to the variation of signal overlap reflecting the composition. Here, principal component analysis (PCA) was applied to the analysis of four components using 20 integral regions [18] indicated in Figure 11.3. Automated analysis was carried out by reading 27 sets of data from SBR and BR samples with 20 integral values, into a spreadsheet program to construct a 20 x 27 independent matrix. Then, a 4 x 27 dependent matrix was created using the composition data obtained according to the equations (11.1). PCA using the partial-least squares regression algorithm was performed on the data set. The result of composition analysis by PCA was almost equivalent within experimental error to that obtained by ordinary 13C-NMR method. Automated sample classification is also achieved by PCA analysis.

408


Table 11.2 Integral regions for PCA analysis of SBR and their structural meaning [18] Integral

13

C-NMR Chemical Resonance Assignment Shift Region (ppm)

Integral in Figure 11.3

1

146.0-145.0

Styrene C1

2

143.2-142.2

Butadiene

vinyl CH

B

3

130.5-130.05

Butadiene

trans CH

C

4

130.05-129.8

Butadiene

trans CH

C

5

129.8-129.5

Butadiene

cis CH

C

6

129.5-129.0

Butadiene

cis CH

C

7

126.2-125.6

Styrene C4

8

114.7-113.7

Butadiene

9

46.2-45.2

Styrene CH

10

44.0-43.0

Butadiene

vinyl CH

11

40.6-39.6

Butadiene

trans CH2 styrene CH

12

38.6-37.5

Butadiene

trans CH2 vinyl CH

13

36.2-35.2

Styrene CH2

14

34.6-33.6

Butadiene

vinyl CH2

15

33.3-32.1

Butadiene

trans CH2

F

16

31.0-30.25

Butadiene

trans CH2 styrene CH2

G

17

30.25-29.5

Butadiene

trans CH2 vinyl CH2

G

18

27.9-26.9

Butadiene

cis CH2

H

19

25.7-25.2

Butadiene

cis CH2 styrene CH2

I

20

25.2-24.7

Butadiene

cis CH2 vinyl CH2

I

A

D vinyl CH2

E

Reproduced with permission from N.G. Walsh, J.K. Hardy and P.L. Rinaldi, Applied Spectroscopy, 1997, 51, 889, Table 1. Copyright 1997, Society for Applied Spectroscopy

409


Figure 11.3 13C-NMR spectrum of SBR including poly(ethylene glycol) as internal standard, observed at 75 MHz. The signals in the region A-I are used to calculate the integral intensity according to the equations (11.1) [18] Reproduced with permission from N.G. Walsh, J.K. Hardy and P.L. Rinaldi, Applied Spectroscopy, 1997, 51, 889, Figure 1. Copyright 1997, Society of Applied Spectroscopy

11.3.2 Analysis of Chemical Composition Distribution using SEC-NMR Copolymers usually have three types of distribution, i.e., the molecular weight distribution, sequence distribution and chemical composition distribution. The chemical composition distribution means the molecular weight dependence of chemical composition. Fractionation techniques, such as solvent fractionation and HPLC, including SEC, are commonly used for the analysis of molecular weight distribution. So-called cross fractionation, using two sets of solvents or special HPLC columns, is applied to the fractionation followed by chemical composition analysis for each fraction. The direct coupling of HPLC (SEC) and NMR spectroscopy will be a powerful tool for the analysis of chemical composition distribution. This is due to the development of high-field NMR instruments and commercially available HPLC-NMR proves, which ensure high sensitivity and resolution for a flowing sample [19, 20]. The applicability of on-line SEC-NMR was demonstrated for the analysis of molecular weight dependence of the chemical composition [21] and tacticity [22]. The quantitative measurement of end-group of a polymer chain has enabled the direct measurement of molecular weight distribution [23]. On-line SEC-NMR has successfully applied to the determination of chemical composition distribution in ethylene-propylene-diene terpolymers (EPDM) by the use of 1H-NMR at 750 MHz [24]. EPDM containing 2-ethylidene-5-norbornene (ENB), as a diene

410

High Resolution NMR of Elastomers component, shows the 1H-NMR signals characteristic of each component. The assignment of signals were made by means of PFG-heteronuclear quantum coherence (HMQC) measurements, using the assignments of 13C-NMR signals of the terpolymer thus far reported [25, 26]. As shown in Figure 11.4, the continuous-flow SEC-NMR gives highresolution signals almost comparable to those recorded under ordinary NMR measurements. It also gives enough signal-to-noise ratio (S/N) for the signals assigned to β+ -CH2 and P-CH3 at 1.25 and 0.83 ppm, respectively. However, the signals due to a small amount of ENB are difficult to use for quantitative measurements in the continuous-

Figure 11.4 1H-NMR spectra of EPDM in deuterochloroform, observed at 750MHz [24]. (Sample contains E 78.9, P 19.1, and ENB 1.1 mol%): (a) Continuous-flow SECNMR measurement, (b) normal measurement using 5 mm o.d. sample tube and (c) stop-and-flow SEC-NMR measurement Reproduced with permission from K. Ute and K. Hatada, International Journal of Polymer Analysis and Characterisation, 1999, 5, 47, Figure 3. Copyright 1999, Academic Press

411

Spectroscopy of Rubbers and Rubbery Materials flow experiments. In the stop-and-flow experiments, NMR scans are accumulated in each point until an enough S/N is obtained for the signals due to ENB component. Here, it is remarkable that the SEC-NMR spectra at 750 MHz instrument improves the S/N ratio by 9.2 times compared with those at 500 MHz one [23]. Figure 11.5 shows a stacked trace plot of the continuous-flow SEC-NMR data for a mixture of EPDM and EP, containing 31.5 mol% P and 0.6 mol% ENB on an average. Cross sections of the stacked plot at 1.25 ppm and 0.83 ppm give the 1H-NMR detected SEC curves for E and P components, respectively, as a function of elution volume, using the following equations: Ip = I(P-CH3)/3 Ie = [3-I(α-CH2) + I(β+ -CH2) - 2Ip]/4 [3-E] = 100Ie /(Ie + Ip) The relationship between the elution volume and chemical composition can be obtained by using a calibration curve of standard samples, such as standard polystyrene. It was clearly shown that the ENB and ethylene contents in this sample decreased with decreasing molecular weight. Here, it should be noted that the elution volume reflects the molecular size including the molecular weight and composition. This measurement was made using deuterochloroform, but it was suggested to use o-dichlorobenzene as eluent at 135 °C to separate mainly by molecular size of EPDM. Temperature gradient fractionation followed by NMR analysis also gives chemical composition distribution for EPDM containing

Figure 11.5 A stacked trace plot of 750MHz continuous-flow SEC-NMR spectra [24] Reproduced with permission from K. Ute and K. Hatada, International Journal of Polymer Analysis and Characterisation, 1999, 5, 47, Figure 6. Copyright 1999, Academic Press

412

High Resolution NMR of Elastomers higher ENB units [27]. This measurement has shown that ENB content is independent of molecular weight, while ethylene content decreases in high molecular weight fraction.

11.3.3 Analysis of Sequence Distribution Sequence distribution studies on several types of rubber by 13C-NMR technique have been reported. Some of the more recent reports include silicone rubbers [28-30], SBR [31], acrylonitrile-butadiene rubber (NBR) [32,33], polyurethane [34,35], polyepichlorohydrin [36], ethylene-norbonene [37] and ethylene-propylene rubber [4, 16, 25, 38-44]. The NMR studies on EPDM have been carried out extensively, because it is one of the important parameters, which control the physical properties of the elastomer. For example, ethylene sequence can influence the crystallisation kinetic and melting behaviour of the rubber [38]. The early definitive work on comonomer sequence distribution was based on mathematical model of ethylene-propylene polymerisation, where propylene added by either primary or secondary insertion [43]. The improvement [39] was made using higher field 13C-NMR to correct several mistakes, which were made in the earlier signal assignments [40-42]. A typical 13C-NMR spectrum of EPM is given in Figure 11.6. The corresponding signal assignments are indicated in Table 11.1.

Figure 11.6 13C-NMR spectrum of EPM in tetrachloroethane-d2 [16] Reproduced with permission from S.D. Martino and M. Kelchtermans, Journal of Applied Polymer Science, 1995, 56, 1781, Figure 1. Copyright 1995, John Wiley & Sons

413

Spectroscopy of Rubbers and Rubbery Materials Later, it was indicated that 13C-NMR signals of the third monomer such as 1,4-hexadiene, ENB and 5,8-dicyclopentadiene might overlap with the main chain EPM carbons, particularly in the Sαα region [43]. The difference in ethylene content obtained by different methods of calculation could be as high as 6 mol% with increasing the ENB content. Using the correction of the third component, the standard deviation dropped from 3% to 0.5%. The technique on sequencing experiments can also be improved by introducing paramagnetic relaxation agent such as Cr(acac)3 to shorten the 13C spin lattice relaxation times. The sensitivity of measurement is further enhanced via NOE by allowing the 1H decoupling during the recycle time because the NOE is expected to be constant for all carbons in the polymer chain. With these improvements, the experimental time was reduced from more than 12 hours to 3-6 hours [43]. An improved method was proposed for the determination of triad distribution, by considering the inherent uncertainties in the quantitative analysis of EP copolymer in EPM from 13C-NMR integral intensities [44]. The 13C-NMR signals due to triad sequence of E and P units were grouped into 8 regions and 8 simultaneous equations were derived to describe six triad sequence distribution numbers. Using a statistical Z-test, the triad concentrations were obtained by solving directly the over-determined set of equations. These results showed smaller standard deviation than those derived from the equations that take into the consideration of systematic errors in the derived triad due to spectral overlap. This was further confirmed by the uncertainty analysis using simulation, assuming a Gaussian error distribution. The results also showed that the uncertainty in a triad concentration is affected by the relative distribution of the other triad concentrations. This effect is larger for the sequence with higher ethylene content. The content of 1,2 (vinyl), cis and trans isomeric units in polybutadiene (PB) can be obtained by 1H-NMR [1]. A better spectral resolution of 13C-NMR permits the sequence analysis to be used to quantify the microstructure of PB. Thus, the assignments for aliphatic [45] and olefinic [46] carbons in triad sequence have been reported. As shown in Figure 11.7, the olefinic carbon resonances were assigned to the triad sequences of 1,2 (vinyl), cis and trans in the 13C-NMR of PB [47]. A good agreement was observed between the experimental peak intensity and calculated one assuming the Bernoullian model. The assignment of peaks in the aliphatic and olefinic regions was further tested using 1,2′,4,4′tetradeuteriopolybutadiene (PB-d4h2) and perdeuterated polybutadiene (PB-d6) in comparison with the nondeuterated PB (PB-h6) [48, 49]. The trans content of deuterated and partially deuterated PB determined from the aliphatic region was higher than that of olefinic region of partially deuterated PB due to signal overlap in the region [49]. Since, only aliphatic region is amenable for quantitative analysis of PB-d6, {2H}13C-INEPT NMR with simultaneous 1H and 2H decoupling during data acquisition was used to characterise the sequence distribution of deuterated PB [50]. By using this technique, spectral simplification is achieved, which is useful for performing quantitative analysis and developing signal assignment. 414


Figure 11.7 Expanded 13C-NMR spectra of the olefinic resonances of polybutadiene, due to cis-1,4 and trans-1,4 centred triads, observed at 50 MHz [47]. The content of 1,2 units is 1% (A), 12% (B), 13% (C), 22% (D), and 49% (E) Reproduced with permission from G. van der Velden, C. Didden, T. Veermans and J. Beulen, Macromolecules, 1987, 20, 1252, Figure 2. Copyright 1987, American Chemical Society

415

Spectroscopy of Rubbers and Rubbery Materials Synthetic cis-1,4 polyisoprene has the structural feature similar to PB. As to the NMR analysis of sequence distribution, however, little work has been done after 1978. Detailed assignment may be possible by using modern NMR techniques and magnetic field higher than 100 MHz for 13C-NMR. The microstructure of polychloroprene, polymerised at +12 to +70 ºC, was analysed using 1H- and 13C-NMR [32]. Signal assignments were made for head-and-tail arrangements of trans-1,4 units, which was the major component of polychloroprene, and for other isomeric units. Polymerisation at high temperature resulted in a slight increase in head-to-head and tail-to-tail linkage of trans-1,4 units as well as the increase in cis-1,4 units. The assignment of 13C-NMR signals in carboxyl-terminated acrylonitrile-butadiene copolymers was done by using one and two-dimensional (1D and 2D) NMR techniques [32]. The fundamental assignment of 1H-NMR signals of NBR done in 1970s was confirmed based on 2D COSY-45 and 2D J-resolved spectra. Detailed assignment of 13 C-NMR signals were successfully made with spectral editing of carbon peaks using DEPT technique as well as C-H COSY and long-range two-dimensional heteronuclear chemical shift correlation (COLOC) experiments, as tabulated in Table 11.3. Block polybutadiene units were found in this polymer, but no block acrylonitrile sequence was detected. The absence of diad sequence of acrylonitrile units in NBR containing acrylonitrile less than 30-33% was also confirmed by 1H-NMR [52] and 13C-NMR [33]. Segmented polyurethane-urea elastomer is a block copolymer composed of soft and hard segments, see Figure 11.8. The composition and length of both segments as well as sequence length distribution are the primary factors to govern the mechanical properties. The mean sequence lengths of polyurethane block and polyurea block, were analysed by 13 C-NMR, based on the assignment of aromatic carbon signals by using model compounds, as depicted in Figure 11.9 [34]. Quantitative analysis indicated that the polymer is almost a random copolymer. A similar analysis of segmented polyurethaneurea elastomer found that two-step polymerisation gave the mean sequence length longer than one-step polymerisation [35]. The combination of lanthanide shift reagents and 2D NMR techniques (COSY and NOESY) was applied to the 1H- and 13C-NMR assignments of ethylene-vinyl acetate

Figure 11.8 Structure of segmented polyurethane-urea

416

High Resolution NMR of Elastomers Table 11.3 Assignments of 13C-NMR peaks of carboxyl-terminated NBR [32] Chemical Shift (δ)

Sequence

Carbon Type

Chemical Shift (δ)

Sequence

Carbon Type

176.40

I

COOH

143.46

VV

-CH=

142.45

VC

-CH=

142.28

VT

-CH=

140.56

VA

-CH=

140.36

VA′

-CH=

140.24

TV′I

-CH=

134.31

TTA

-CH=

134.07

TTAT

-CH=

133.06

TV′I

-CH=

133.01

TV′A

-CH=

132.41

ATA

-CH=

131.73

TTTA

-CH=

131.27

ATTA

-CH=

131.17

A′TTA

-CH=

130.15

VTT

-CH=

129.93

CTT

-CH=

129.80

TTT

-CH=

129.59

TTV

-CH=

129.40

TVT

-CH=

129.22

TCT

-CH=

129.05

TCV

-CH=

128.20

TTTA

-CH=

127.85

ATTA

-CH=

126.81

CyCA

-CH=

126.53

ATA

-CH=

126.40

ATAT

-CH=

124.85

TTA

-CH=

124.75

TTAT

-CH=

124.66

TTA

-CH=

124.25

ATV′

-CH=

124.02

ATV′

-CH=

123.11

CyCA, I

-CH=,CN

122.20

CyCA

CN

122.05

AC

CN

121.64

TAT

CN

121.51

AT

CN

120.26

AA

CN

117.15

AV

=CH2

116.44

TV′I

=CH2

115.71

AV′

=CH2

114.31

VT

=CH2

114.09

VC

=CH2

43.31

TVT

CH

42.44

TV′I

CH, e or ta

41.28

TV′I

CH, e or ta

42.03

TV′I

=CH2

41.88

VA

CH, e or ta

41.03

VA

CH, e or ta

40.43

TVI

CH

417

Spectroscopy of Rubbers and Rubbery Materials Table 11.3 Continued... Chemical Shift (δ)

Sequence

Carbon Type

Chemical Shift (δ)

Sequence

Carbon Type

38.40

AVT

CH 2

37.96

VT

CH2

37.20

TV′I

CH2

36.34

TA′I

CH2

36.09

I

C

35.45

TVV

CH 2

34.86

TA′

CH2

34.70

AAT

CH2

33.86

TV

CH2

33.21

I

CH2

33.11

I

CH2

32.49

TTT, VC

CH2

32.36

TTA

CH2

32.14

TTTA

CH2

32.05

TTAA

CH2

31.34

AT

CH2

31.21

TAT

CH2

31.05

TAT + AT

CH

30.85

A′VT

CH 2

30.77

AV′T

CH

30.11

I

CH 2

29.95

TV

CH2

29.78

TA

CH2

29.54

AC

CH2

29.41

CI

CH2

29.10

VAT

CH, e or ta

28.10

A′I

CH, e or ta

27.94

CyCA

CH2

27.28

CC

CH2

27.16

CT

CH2

26.94

CCA

CH2

26.78

CC

CH2

25.10

CyCA

CH2

24.95

CI

CH2

24.85

CVA

CH2

24.66

CV

CH2

24.60

CA

CH2

24.41

CyCA

CH

23.28

TV′I

CH 3

22.78

CyCA

CH2

22.15

TVI

CH3

A: acrylonitrile unit, B: butadiene unit, V: 1,2-butadiene unit, C: cis-1,4 butadiene unit, primed units: reversed arrangement, T: trans-1,4 butadiene unit, CyCA: 2-cyclohexene1-carbonitrile and I: 4,cyanopentanoic acid end group a Erythro or threo defined by a Newman projection of this carbon and the other nearest asymmetric centre Reproduced with permission from T. Fang, Macromolecules, 1990, 23, 2145, Table 2. Copyright 1990, American Chemical Society

418


Figure 11.9 13C-NMR spectrum of segmented polyurethane-urea, observed at 75.5 MHz [34] Reproduced with permission from A. Kajii and M. Murano, Polymer Journal (Japan), 1990, 22, 1065, Figure 3. Copyright 1990, The Society of Polymer Science, Japan

copolymer [53]. The addition of shift reagents causes a number of resonances to emerge and shift downfield from the large resonances centered at 1.26 ppm, as summarised in Table 11.4. The complete 13C-NMR interpretation was obtained for the sequences.

11.3.4 Analysis of Chemically Modified Structure and Graft Polymers Chemical modification of rubber such as hydrogenation, chlorination, epoxidation, cyclisation, addition of maleic anhydride, is a popular technique to give a desirable

419

Spectroscopy of Rubbers and Rubbery Materials Table 11.4 13C- and 1H-NMR assignments in EVAc copolymers [53] 1

13

H

Shift [Yb(fod)3]a

C

No.

Shift [ρ=0]

1

1.26

ca. 1.26

Sγγ+Sγδ+Sδδ

2

1.33

1.60

3

1.40

4

Sequence

Sequence

Shift

Acetate

21.0

S βδ

Sββ

21.5

ca. 1.70

Sββ

Sβδ

25.7

1.50

1.93

Sαγ+Sαδ

Sγγ+Sγδ

29.8

5

1.73

2.08

S αα

Sδδ

30.0

6

1.92

2.58

Acetate

Sαδ

34.7

2.74

Sαγ

35.2

2.90

Sαα

39.6

Tββ

ca. 68

Tβδ

71

Tδδ

73

7

5.15

5.35

Tββ (VVV)

8

5.03

5.75

Tβδ (VVE)

9

4.90

5.90 Tδδ (VEVEV) 6.00 Tδδ (VEVEE) 6.30

Tδδ (EEVEE)

a Shift observed with Yb(fod)3 added, reagent/substrate molar ratio ρ = 0.00235 Reproduced with permission from H.N. Cheng and G.H. Lee, Polymer Bulletin, 1988, 19, 89, Table 1. Copyright 1988, Springer-Verlag GmbH & Co KG

property. Structural characterisation of the chemically modified rubbers is done to decide the extent of reaction, sequence distribution of modified units and structural change by side reactions. Acid-catalysed cyclisation of polyisoprene and polybutadiene forms resin-like polymer containing cyclised segments consisting of one or more fused rings as shown in Figure 11.10. The average number of rings per segment, termed cyclicity, degree of cyclisation and segment end-groups have been analysed by the use of 1H-NMR and 13C-NMR. The cyclicity was presumed to be from three to five in extensively cyclised polyisoprene with TiCl4/CCl3CO2H catalyst [54].

420


Figure 11.10 Cyclised polyisoprene with tetracyclic segment Hydrogenation of NBR has been developed to improve thermal stability. Perfect hydrogenation of the double bonds in butadiene unit, without any hydrogenation of the acrylonitrile units, using RhClH2(P(C6H5)3)2 catalyst was confirmed with 1H- and 13 C-NMR [55]. In the 13C-NMR spectrum of hydrogenated NBR, CN carbon atom showed splitting reflecting the pentad sequences of acrylonitrile and ethylene units [56]. The 1H-NMR analysis showed clearly the progress of ring-opening reaction of the epoxide group to form diols in partially epoxidised natural rubber [57]. Formation of fluoroesters of cis-polyisoprene (Guayule rubber) by ring opening with fluoroacids was confirmed with 1H- and 13C-NMR [58]. The 13C-NMR signals of epoxidised low molecular weight polybutadiene have been assigned to diad sequences of epoxide units based on the calculation of chemical shifts by empirical rule. The cis-1,4 butadiene units epoxidised preferentially and no reaction was observed for 1,2 units [59]. The functionalisation of polyethylene, polypropylene and EPM with maleic anhydride (MA) has been done to give compatibilisation and enhanced adhesion to polar materials. The free-radical initiated grafting of 13C-enriched [2,3-13C2]MA on to polyethylene (PE), polypropylene (PP) and EPR was studied using noise-decoupled and 1D INADEQUATE (incredible natural abundance double-quantum transfer experiment) 13C-NMR [60,61]. The result indicated the formation of following graft structures shwon in Figure 11.11, where MA grafts were in the form of single succinic anhydride part in the case of EPR and PP. In EPM with high propylene content, a chain scission reaction could occur, yielding a ring attached to the chain-ends via a double bond. The structure of curing points has been studied using model compounds of rubber.

Figure 11.11 Presumed structure of polyethylene and polypropylene modified with maleic anhydride

421

Spectroscopy of Rubbers and Rubbery Materials Vulcanisation of EPDM with sulfur systems was studied by 1H-NMR using ethylidene norbornane (ENBH) as a model of ENB [62]. The use of ENBH was also effective to elucidate the curing reaction of EPDM with phenol-formaldehyde resin [63]. Similarly, halogenation reaction of IIR was studied by 1H-NMR using 2,2,4,8,8-pentamethyl-4nonene as a model [64].

11.4 Analysis of End-groups and Branching

11.4.1 Assignment of Small Signals Structural characterisation of chain end-groups provides information on the mechanism of polymer formation and degree of polymerisation. The importance of functional endgroups on the performance of elastomers has been recognised recently. For example, chemical modification of end-groups has led to the improvement of dynamic properties of anionically polymerised SBR [65,66]. Recent progress of NMR instruments enabled the quantitative analyses of end-groups in high molecular weight polymers with less difficulty. This is because the limit of detection is sufficiently high to analyse the number average molecular weight of a polymer based on the end-group analysis, if the polymer molecule contains a known amount of end-groups per chain. For instance, the endgroup analysis makes it possible to determine the M n value higher than 50 x 104 for PMMA, by the use of 1H-NMR at 500 MHz [11]. Chemical modification of chain endgroups is an ordinary technique to obtain a special functional group. The extent of reaction can be evaluated by the functionality of polymer chain, i.e., the average number of functional group per chain. Long-chain branching in polymer chain could also give rise to small NMR signals. The signal assignment for branch-points is sometimes more difficult than that of end-groups, because of difficulty in the preparation of low molecular weight model compounds. One of the major challenges would be assignment of small signals. However, the use of modern NMR instruments at high magnetic field has enabled one to apply various NMR techniques such as 2D NMR and DEPT, for the assignment of very small signals. Other difficult tasks may include the purification of polymer samples. Sometimes it is important to distinguish between the signals due to end-groups or branch-points in a polymer chain and signals from impurity, especially in the case of naturally occurring rubbers. It is preferable to remove impurities, which could give rise to signals near the resonance in question. This is to ensure that the signal intensity of impurities is much smaller than that of chain-ends, abnormal groups or branch-points, in order to get unequivocal assignment.

422


11.4.2 Functionality of Telechelic Diene Polymers The chain-ends of anionically polymerised so-called “living” polybutadiene and polyisoprene can be modified to form hydroxyl or carboxyl group at the final stage of polymerisation. The functionality of telechelic diene polymers is a very important factor to establish a uniform network structure by a coupling reaction. A simple chemical titration or NMR analysis can be applied to the determination of functionality, because telechelic rubbers are normally of low molecular weight. The functionality of hydroxy-telechelic polybutadiene prepared by the reaction of ethylene oxide was reported to be 2.03 ± 0.1, by measurement at 100 MHz in the early stage [67]. Here, the functionality was determined according to the following equation, using the relative intensity of signals in Figure 11.12.

Here, x and y stand for the weight fraction of 1,2 and 1,4 units, respectively. On the other hand, commercial hydroxy-telechelic polybutadiene obtained by radical

Figure 11.12 1H-NMR spectrum of hydroxy-telechelic polybutadiene, observed at 100 MHz [67] Reproduced with permission from Y. Camberlin, J.P. Pasucault and Q.T. Pham, Die Makromolekulare Chemie, 1979, 180, 397, Figure 1. Copyright 1979, Wiley-VCH

423

Spectroscopy of Rubbers and Rubbery Materials polymerisation had a distribution of functionality. The Fn value increased from 1.9 to 3.5 (average Fn =2.1) with increasing the molecular weight, as it was shown by the analysis of fractions with SEC fractionation [68]. The functionality of hydroxy-telechelic polybutadiene and polyisoprene was determined by 1H-NMR after modification with phenyl isocyanate or naphthyl isocyanate. Liquid telechelic polybutadiene and polyisoprene, prepared by choosing appropriate reaction conditions, showed Fn value of about 1.9, which agreed with the values determined by titration method [69]. Amino-terminated telechelic polybutadiene was prepared by LiAlH4 reduction of amidino end-group in polybutadiene, which was polymerised by a water-soluble initiator, 2,2′azobis(amidinopropane)dihydrochloride. The structure was analysed by 1H- and 13C-NMR, but functionality of 2.0 was obtained by a titration method [70]. Synthesis of ω-epoxyfunctionalised polyisoprene was carried out by the reaction of 2-bromoethyloxirane with living polymer that was initiated with sec-butyllithium. The functionality of the resulting polyisoprene was 1.04 by 1H-NMR and 1.00 by thin layer chromatography detected with flame ionisation detection [71]. The terminal hydroxyl groups in hydroxy-telechelic polybutadiene are assumed to be predominantly primary alcohols linked to cis-1,4, trans-1,4, and 1,2 butadiene units [72]. Detailed structure of the butadiene or isoprene units at both terminals was studied by the use of 2H-NMR [73], as shown in Section 6.1

11.4.3 Structure of Terminal Groups The 13C-NMR of low molecular weight ethylene-propylene rubber (EP) exhibited 4 olefinic and 33 aliphatic signals. These signals were assigned to the diad and triad sequences of E and P units including initiating and terminating units. Empirical additive rules have been adopted to estimate the chemical shifts for the structures expected from the proposed polymerisation mechanism [42]. A complete triad sequence distribution, including all of the possible end-groups, was obtained for EP with metallocene catalyst as shown in Figure 11.13 [74]. A statistical approach was used, based on a first-order Markovian distribution, which indicated that initiating units were almost exclusively propyl groups. Mixing and extrusion techniques are usually applied to elastomers for processing and improvement of physical properties. These mechanical processes normally lead to the occurrence of radical reactions such as chain scission, chain coupling, and crosslinking. The 13C-NMR analysis of EP showed that the shearing of polymer resulted in hydrogen abstraction. This was followed by disproportionation reaction to form olefins [75].

424


(a)

(b)

Figure 11.13 Structures and 13C-NMR signal assignments for: (a) olefinic end-groups, and (b) aliphatic end-groups. P denotes the polymer chain and polymerisation occurs from left to right [74] Reproduced with permission from J.C. Randall and S.P. Rucker, Macromolecules, 1994, 27, 2120, Figures 3 and 4. Copyright 1994, American Chemical Society

Detailed assignment of signals was carried out for various types of end- group of EP with molecular weigh ca. 1.7 x 103, obtained after high temperature extrusion [76]. In this case, 2D NMR measurements at 300 Hz were able to be applied due to the low molecular weight of the samples. A mechanism involving hydrogen abstraction followed by disproportionation is presumed to form different types of olefins. The 1H-NMR signals in the olefinic region were assigned based on information from 1H-13C 2D correlation measurement. The molecular weight was estimated based on the relative intensity of olefinic proton signals from the endgroups. A reasonable agreement was observed between the molecular weight by 1 H-NMR and by SEC for EP samples decomposed by shearing at various temperatures [76]. The concentration of end-groups increased in the order of vinylidene>vinyl>vinylene for all samples prepared at extruder shearing temperature from 260° to 450 °C.

425

Spectroscopy of Rubbers and Rubbery Materials The mechanism of γ-ray irradiation-induced scission of polyisobutylene was studied, based on the structural characterisation of end-groups by 13C-NMR as well as GC, GC/ MS, and SEC [77]. The assignments of signals were made by comparison with those from model compounds and predictions based on empirical rules. Quantitative 13C-NMR measurements of chain-ends allowed the determination of radiation yield of products and of chain scission. Carboxyl-terminated butadiene-acrylonitrile copolymer, mentioned above, was found to have 4-cyanopentanoic acid end-group originating from 4,4,7-azobis(4-cyanopentanoic acid) initiator [32].

11.4.4 Structure of Branch-points and Coupling Points The incorporation of a special functional group at the chain-end of rubber molecules was recognised as a way to improve dynamic properties of vulcanised rubber with carbon black, such as resilience, heat build-up, and tan δ [65]. Some coupling agents, used to combine chain-ends [66] or chains to carbon black [78], are also known to play a similar role. Living styrene-butadiene copolymer having butadienyl-Li end-groups was prepared with n-butyllithium/THF initiator by adding a small amount of butadiene at the final stage of polymerisation. The resulting living copolymer was coupled with tin tetrachloride [66]. The presence of terminal butadienyl carbon linked to tin was presumed based on Sn-C-C=C- signals around 11 and 14 ppm in the 13C-NMR spectrum, which were also observed in model compounds. The resulting rubber showed a bimodal molecular weight distribution, suggesting the presence of unreacted chains, although quantitative measurements were not presented. It is remarkable that the coupling of butadienyl-tin is more effective than that of styryl-tin for the improvement of dynamic properties. A similar modification at the chain-end of neodymium catalysed living cis-1,4 polybutadiene was also attempted [79,80]. Butyl rubber is a copolymer of isobutylene with 1-3 mol% isoprene. The isoprene units are presumed to be predominantly in the trans-1,4 addition and in head-to-tail arrangement [81,82]. However, a very small doublet signal appears at 4.93 ppm in 1HNMR, in addition to a triplet signal at 5.05 ppm, which was assigned to the olefinic proton of isoprene units in 1,4 addition, as shown in Figure 11.14. The doublet signal was tentatively assigned to the isoprene unit in 1,2 or 4,3 addition [14,83]. Detailed assignment of this signal was made by using 2D methods at 500 MHz, such as 1H-1H COSY and hypercomplex phase-sensitive inversely detected heteronuclear multiplequantum coherence (HMQC) [84]. These experiments indicated that the signal at 4.93 ppm corresponds to =CH coupled with a signal at 2.44 ppm having equal intensity and 1.72 ppm having 3 times in the intensity. As a model, isobutylene-d 8 was

426


Figure 11.14 1H-NMR spectrum of (a) commercial butyl rubber and (b) polyisobutylene, observed at 500 MHz [84] Reproduced with permission from J.L. White, T.D. Shaffer, C.J. Ruff and J.P. Cross, Macromolecules, 1995, 28, 3290, Figure 1. Copyright 1995, American Chemical Society

copolymerised with isoprene. The aliphatic methine signal at 2.44 ppm appeared as a sharp doublet with the same splitting as that of 4.93 ppm signal. These findings indicate that the signals are derived from the isoprene unit in Figure 11.15, consisting of a branchpoint. The presence of long-chain branching was supported by SEC and rheological measurements.

Figure 11.15 Presumed structure of the branch point of butyl rubber

427

Spectroscopy of Rubbers and Rubbery Materials The 1H-, 13C- and 2D NMR techniques were applied to the analysis of short chain branching of ethylene copolymers synthesised at high pressure by free-radical initiator [85]. The level of short chain branching (branches of length C5 or shorter) increases from about 6 to 20 branches per 1000 CH2 for ethylene homopolymer with increasing the polymerisation temperature from 165 °C to 270 °C. By comparison with the spectrum of ethylene homopolymer, the assignment of 13C-NMR signals due to branches was made for ethylene copolymers, such as n-butyl acrylate, methyl acrylate, vinyl acetate, n-butyl methacrylate, acrylic acid, and methacrylic acid. Confirmation of the assignments has been obtained in a 2D HSQC-TOCSY (heteronuclear single quantum correlation-total correlation spectroscopy) experiments. Ethylene-vinyl acetate copolymer (EVA) showed two additional signals in the ethyl region of the 13C-NMR spectrum, as shown in Figure 11.16. In contrast to the acrylate copolymers, branching structure of EVA is more similar to ethylene homopolymer as summarised in Table 11.5.

Figure 11.16 13C-NMR signals due to short chain branches in ethylene-vinyl acetate copolymers, observed at 100 MHz: (a) ethylene homopolymer (195 °C), (b)14 wt% VAc (200 °C), (c) 22 wt % VAc (200 °C) and (d) 38 wt % VAc (200 °C) [85]. (The notation of carbons is given in Table 11.4) Reproduced with permission from E.F. McCord, W.H. Shaw Jr. and R.A. Hutchinson, Macromolecules, 1997, 30, 246, Figure 10. Copyright 1997, American Chemical Society

428


Table 11.5 Short chain branch (SCB) structures in ethylene-vinyl acetate copolymers from 13C-NMR analysis a [85] Trct (°C)

wpVAc

VActot

1B2

2B4

1A3B4+

2B5+

SCBtot

165

0.000

0

0.9

3.3

0

3.2

7.4

165

0.141

26

1.1

3.7

0.7

3.3

9.0

165

0.235

48

0.7

3.4

0.7

3.3

8.3

165

0.386

93

1.4

2.8

1.6

4.0

10.0

195

0.000

0

1.7

4.3

0

3.7

9.7

200

0.140

26

1.9

5.7

0.7

4.9

13.8

200

0.219

44

1.4

4.8

1.2

4.9

12.8

200

0.230

46

1.5

5.3

1.2

5.1

13.4

200

0.381

91

0.9

3.7

2.2

4.9

12.0

a

All numbers (columns 3-8) per 1000 CH2 units The second carbon from the end of a butyl branch is 2B4. xBy+ refers to branches of length y and longer and includes end-of-chain (EOC) contributions. xAQBy indicates that the comonomer unit located off the backbone carbon at the branch point, making it quaternary Reproduced with permission from E.F. McCord, W.H. Shaw and R.A. Hutchinson, Macromolecules, 1997, 30, 246, Table 3. Copyright 1997, American Chemical Society

11.5 Structural of Naturally Occurring Polyisoprenes

11.5.1 Structure of Natural cis- and trans-Polyisoprenes Naturally occurring polyisoprenes are composed of isoprene-units in the cis or trans configuration. Natural rubber from H. brasiliensis (NR) and Gutta percha from P. gutta are typical cis-1,4 and trans-1,4 polyisoprene, respectively. On the other hand, Chicle from A. zapota contains both cis-1,4 and trans-1,4 polyisoprenes. NR has been presumed to be a pure cis-1,4-polyisoprene, containing no detectable amount of trans-1,4 and 3,4 units as indicated by 1H-NMR at 100 MHz [86]. Natural isoprene oligomers consisting of 9-23 isoprene-units, referred to as polyprenol, can be a good model for structural

429

Spectroscopy of Rubbers and Rubbery Materials analysis of naturally occurring polyisoprenes. Polyprenol is classified into three categories, i.e., all-trans (Solanesol-type), two-trans and poly-cis (Betulaprenol-type), and three-trans and poly-cis (Ficaprenol-type). However, a polyprenol of all-cis configuration has not been reported [87]. The sequence distribution of both trans and cis isoprene-units in polyprenol has been analysed by 13C-NMR, using acyclic terpenes of all the arrangements of both trans and cis isoprene-units, corresponding to isoprene dimer to tetramer. The chemical shifts of signals reflecting the sequence distribution are independent of the degree of polymerisation for polyprenols longer than 10-mer [88-90], as shown in Figure 11.17. Here, the dimethylallyl group and hydroxylated isoprene-unit are referred to as ω and α terminals, respectively, and the carbon atoms in isoprene-unit are numbered as indicated in Figure 11.17. The absence of cis-trans linkage in all the polyprenols and the presence of only one trans-cis linkage per chain indicate that the trans-cis polyprenols are di-block copolymer, aligned in order of ω-(trans)2 or 3-(cis)n-OH. Solanesol is a homologue of alltrans acyclic terpenes.

Figure 11.17 Relationship between length of isoprene-units and chemical shifts of 13CNMR signals reflecting the arrangement of isoprene-units, observed at 50 MHz [88] Reproduced with permission from Y. Tanaka in NMR and Macromolecules, ACS Symposium Series No. 247, Ed., J.C. Randall, 1984, ACS, Washington, DC, p241, Figure 3. Copyright 1984, American Chemical Society

430

High Resolution NMR of Elastomers Polyisoprene containing both cis and trans isoprene-units can be prepared by UV irradiation in the presence of thiobenzoic acid [91]. The 13C-NMR study has indicated that cis-trans isomerised polyisoprene is a random copolymer of both isomeric units [92]. Trans-1,4 polyisoprene from Chicle, separated by crystallisation from hexane solution, exhibits the same 13C-NMR signals as those of solanesol, including both terminal groups [93]. This indicates that trans polyisoprene in Chicle is a high molecular weight homologue of solanesol, consisting of about 100 trans-1,4 isoprene-units. Cis polyisoprene in Chicle showed no signal due to the cis-trans linkage, suggesting that Chicle polyisoprene is a mixture of both pure trans-1,4 and cis-1,4 polyisoprenes. Gutta percha and Balata contain high molecular weight ( M n = 1.4-1.7 x 105) trans-1,4 polyisoprene [94,95]. Both shows 13C-NMR signals due to the dimethylallyl group and hydroxylated terminal group. It is remarkable that the diphosphate terminal group was observed in the low molecular weight trans-1,4 polyisoprene from Eucommia ulmoides [96]. The C-4 methylene protons of terminal trans isoprene-unit showed a splitting by coupling from both =CH and 31P in the 1H-NMR spectrum. The splitting pattern is the same as that observed for solanesyl diphosphate. The presence of dimethylallyl group (ω-terminal) and trans isoprene-units in cis-1,4 polyisoprene was clearly observed in the 13C-NMR spectrum of rubber from Lactarius volemus, a rubber-producing mushroom [97,98]. The structure of mushroom rubber was determined to be a high molecular weight homologue of two-trans polyprenol, based on the signal assignments mentioned above, as depicted in Figure 11.18 (1) [89, 93, 95]. The presence of trans isoprene-units and both terminal groups was also confirmed in the low molecular weight rubbers from the leaves of Goldenrod (Solidago altissima) and Sunflower (Helianthus annuus) [94,99]. The arrangement of trans isoprene-units is observed more in detail from the 13C-NMR signal splitting of C-1 methylene carbon [94, 97], as shown in Figure 11.19. The mushroom rubber shows two peaks due to the trans isoprene-unit in dimethylallyltrans and trans-trans linkages similar to those of the two-trans polyprenol. On the other hand, the rubbers from Goldenrod and Sunflower exhibit three peaks similar to the overlap of peaks of two-trans and three-trans polyprenols. This suggests the presence

Figure 11.18 The structure of rubber from Lactarius volemus

431


Figure 11.19 Splittings of the C-1 methylene carbon signal of: (a) rubber from L. volemus, (b) three-trans Ficaprenol-11, (c) two-trans Betulaprenol-18, (d) rubber from Sunflower, (e) rubber from Goldenrod, (f) three-trans Ficaprenol-12 and (g) two-trans Betulapreno-16. (13C-NMR measurements at 100 MHz for (a) to (c) at room temperature for (d) to (f) at 50 °C)

of two types of rubber chains containing two- or three-trans isoprene-units as a mixture, which could be derived from different initiating species of rubber formation such as trans,trans-farnesyl diphosphate (FDP) and trans,trans,trans-geranylgeranyl diphosphate (GGDP) [100].

432


11.5.2 Structure of Natural Rubber

11.5.2.1 Trans Isoprene-units in Natural Rubber Natural rubber from H. brasiliensis showed clear 13C-NMR signals due to trans isopreneunits in the trans-trans linkage, the relative intensity of which decreases with increasing the molecular weight of fraction, as shown in Figure 11.20. These trans isoprene-units were presumed to be derived from an initiating species of rubber formation and not isomerisation of cis isoprene-units, because the trans isoprene-units were in the transtrans linkage and not in the cis-trans linkage [101]. The 1H-NMR spectrum of polyprenols at 500 MHz gives isolated signals for the methyl protons of trans isoprene-units in transtrans and trans-trans-trans sequences [102]. The methyl proton signals of NR shows striking resemblance to two-trans polyprenol, although the signal from dimethylallyl group itself is not detected as shown in Figure 11.21. Accordingly, the structure of NR was postulated to be a kind of two-trans and poly-cis rubber as depicted in Figure 11.22, where ω- and α-ends are unidentified terminal group [103]. Wild rubbers, occurring as latex from higher plants such as Sorvinha (Coma utilis), Indian Laurel (Ficus retusa) and Jelutong (Dyera costulata), show no signal due to trans isoprene-units as well as both terminal groups in the 13C-NMR spectrum [94, 104]. The absence of these groups was presumed to be due to a certain chemical or biochemical oxidative degradation of the rubber molecule during storage in the tree. A similar oxidative degradation of rubber in the tree was observed for NR from firstly tapped Hevea tree, i.e., virgin tree. The rubber from a virgin tree showed no signal due to trans isopreneunits [105,106]. NR contains about 2% proteins and 1% lipids. Deproteinisation of NR with a proteolytic enzyme in latex state reduced the gel content, indicating the decomposition of branchpoints consisting of proteins [107,108]. Transesterification of deproteinised NR (DPNR) with sodium methoxide decomposes fatty acid ester and phosphoric ester linkages in NR to form linear rubber chains [107-109]. The amount of trans isoprene-units per rubber chain was determined from the ratio of cis and trans isoprene-units by 13C-NMR, by considering the number-average molecular weight obtained by osmometry. The observed number of trans isoprene-units in transesterified DPNR was about two independent of the molecular weight [103,109]. The presence of exactly two trans isoprene-units was also confirmed by comparing the 1H-NMR spectrum with those of two-trans and threetrans polyprenols as mentioned above (cf. Figure 11.21).

433


Figure 11.20 13C-NMR spectra of fractionated deproteinized NR from fresh field latex with M n of (a) 2.7 x 105, (b) 0.93 x 105 and (c) 0.30 x 105, observed at 100 MHz [103] Reproduced with permission from Y. Tanaka, A.H. Eng, N. Ohya, N. Nishiyama, J. Tangpakdee, S. Kawahara and R. Wititsuwannakul, Phytochemistry, 1996, 41, 1501. Copyright 1996, Elsevier Science

434


Figure 11.21 1H-NMR spectra of (a) low molecular weight fraction of deproteinized NR (b) two-trans Betulapreno-16 and (c) three-trans Ficaprenol-12, observed at 400 MHz [103]

Figure 11.22 Presumed structure of natural rubber

435


11.5.2.2 Structure of Terminal Groups The absence of dimethylallyl-group in NR indicates that the initiating species for rubber formation in Hevea tree is not FDP, but FDP modified at the dimethylallyl-group, which is abbreviated here as ω′ [103,109,110]. This was confirmed by 13C-NMR analysis of in vitro polymerised rubber by incubation of the bottom fraction of fresh latex and isopentenyl diphosphate (IDP) [111]. The newly synthesised in vitro rubber formed in the presence of FDP and IDP showed the dimethylallyl group derived from FDP. On the other hand, no dimethylallyl group was detected in the in vivo rubber prepared without the addition of FDP [112]. It was disclosed that NR contained about two long-chain fatty acids per rubber chain linked to rubber chain, which were able to remove after transesterification or saponification of NR [107,113]. Most of NR molecules were presumed to be terminated with a phospholipid, taking into the account of the fact that most of rubber particles in fresh latex have no activity for chain elongation reaction [113]. Based on these findings NR was presumed to contain peptide group and phospholipid or phosphate group at the initiating and terminating chain-ends, respectively. Both functional terminal groups can form branchpoints, which are able to decompose by deproteinisation and transesterification (or saponification) [109]. Outstanding properties of NR was presumed to be derived from the long-chain fatty acids included in phospholipid at the α-terminal [109].

11.6 Application of High-Resolution NMR

11.6.1 Multinuclear High-resolution NMR

11.6.1.1 2H-NMR Multinuclear high-resolution NMR other than 1H, 19F and 13C has a versatile application for the structural analysis of polymer, because some nuclei give information about the structure of polymer chain liked to terminal groups, coupling points, and mechanisms of initiation and termination reactions. The structure of chain-end group of living polybutadiene and polyisoprene, quenched with deuteromethanol, was analysed by 2HNMR. The following terminal groups of butadiene and isoprene oligomers, synthesised with ethyllithium in hexane, were confirmed by 2H-NMR [73] (see Figure 11.23). A catalyst system composed of neodymium chloride/trialkyl aluminium is presumed to give living polymer of butadiene containing cis-1,4 units higher than 98%. This living

436


Figure 11.23 Presumed structure of chain-end groups of polybutadiene and polyisoprene initiated with ethyllithium polymerisation of isoprene is reported to accompany side reactions such as chain transfer to monomer. The polyisoprenes quenched with deuteromethanol showed signals due to 3,4 (36-45%) in addition to cis-1,4 terminal units [73], as shown in Figure 11.24. In polybutadiene, trans-1,4 terminal units were predominant, but the formation of 1,2 terminal group was also observed. The initial unit of polyisoprene polymerised with neodymium catalyst in the presence of DAl(i-Bu)2, which is known to act as chain transfer agent, was predominantly trans-1,4 units. It was explained that the principal site of the chain transfer reaction was the Al-D bond. The 2H-NMR was applied to the analysis of cross-linking and branching structure in copolymers composed of butyl acrylate, methyl methacrylate, and small amounts of methacrylic acid and allyl-d5 methacrylate (ALMA-d5) to generate crosslinks. The solution state 2H-NMR spectrum of the copolymer showed clear signals due to the ALMA-d5 units in branch-points. The allyl group in ALMA-d5 was presumed to react with a growing radical chain predominantly by the reaction of the vinyl group at the side chain as shown in Figure 11.25 [114].

11.6.1.2 7Li-NMR 7

Li-NMR has been widely applied to the structural characterisation of lithium initiators for anionic polymerisation. Anionic polymerisation of diene monomer with lithium

437


Figure 11.24 2H-NMR spectra of polyisoprenes obtained with Nd/Al catalyst and quenched by CH3OD, observed at 76.77 MHz [73] Reproduced with permission from K.D. Skuratov, M.I. Lobach, A.N. Shibaeva, L.A. Chunyaeva, T.V. Erokhina, L.V. Osetrova and V.A. Kormer, Polymer, 1992, 33, 5197, Figure 1. Copyright 1992, Elsevier Science

438


Figure 11.25 Formation of branch point in polybutylacrylate initiators gives living polymer. Bifunctional lithium compounds are very important initiators for the preparation of telechelic diene polymers and triblock copolymers. Diene polymers composed of 1,4 units are obtained through organolithium compounds in nonpolar solvents. However, the aggregation or association of both initiators and living chains sometime causes the formation of a broad molecular weight distribution. Initiator systems can be modified by the addition of a special amine or ether and so-called seeding reactions to form polydienyllithium by the addition of small amounts of a diene monomer. 7 Li-NMR gives information on the association, i.e., the most associated species show a signal at the lowest chemical shift. A dilithium initiator, 1,3-phenylene-bis(3-methyl-1-[methylphenylene]-pentylidene) dilithium produced triblock copolymers of styrene-butadiene-styrene (SBS) and styreneisoprene-styrene (SIS) in hydrocarbon solvent. However, resulting copolymers showed broad and bimodal molecular weight distribution [115]. 7Li-NMR indicated that this compound exists as a complex mixture of aggregated species, showing wide chemical shift ranges, while a monolithium compound ((1,1-diphenyl-3-methylpentylidene)lithium) and sec-butyllithium gave a sharp signal. Here, the monolithium compound was presumed to be in dimeric aggregation. The addition of equivalent amounts of N,N,N′,N′′,N′′pentamethyldiethylenetriamine (PMDETA) to the dilithium compound changed these signals to a sharp signal shifted downfield, showing the deaggregation. The addition of tert-butyllithium to 1,2-bis(isopropenyl-4-phenyl)ethane in hexane leads to the formation of a dilithium compound. The bifunctional initiator in hydrocarbon solvent provided diene polymers with unimodal molecular weight distribution. The molecular weight of resulting polymer was the same as that of theoretical one [116]. 7Li-NMR of the resulting α, ω-diithio-polyisoprenes showed a major signal at 0.03 ppm, which was almost the same as the signal of monolithio-polyisoprene, and small signals at -1.64, -1.85 and -1.92 ppm, as given in Figure 11.26. The major signal was assigned to dimeric aggregation and small signals tentatively to cyclic aggregates involving cyclic and linear species involving four lithium atoms. The aggregate species were presumed to be dependent on the chain length. It was also observed that α,ω-diithio-polyisoprenes of high degree of polymerisation exhibited 7Li-NMR similar to those of monofunctional ones [117].

439


Figure 11.26 7Li-NMR spectrum of α,ω-dilithio-polyisoprenes (DPn = 3), 4 x 10-4 M in hexane/C6D6 at 7 °C, observed at 194.3 MHz [116] Reproduced with permission from J.M. Boutillier, J.C. Favier and P. Hemery, Polymer, 1996, 37, 5197, Figure 4. Copyright 1996, Elsevier Science

11.6.1.3 29Si-NMR 29

Si-NMR has been used to determine chain and block length of oligo- and polysiloxanes. The 29Si chemical shift ranges for some structural units in polysiloxanes have been studied [118]. Poly(dimethylsiloxane) and poly(dimethylsiloxane-co-methyl(trifloropropyl)siloxane) were analysed using 13C- and 29Si-NMR [119]. A relatively long spin-lattice relaxation time (T1) and negative NOE were overcome by the addition of a shiftless relation reagent

440

High Resolution NMR of Elastomers such as Fe(acac)3 and by using inverse gated decoupling technique, respectively, to make quantitative 29Si-NMR measurement. The sequence length in poly(dimethylsiloxane)/ poly(tetramethyl-p-silphenylene-siloxane) (PDMS-PTMPSS) copolymers have been analysed by 29Si-NMR [29]. The 29Si-NMR spectrum of copolymer exhibits five signals reflecting the sequence distribution, as shown in Figure 11.27. The assignments were made based on known 29Si chemical shifts. It was shown that redistribution of siloxane bonds proceeded during polymerisation by the presence of water. Sequencing of siloxane units was also investigated for anionic ring-opening polymerisation of 1,1-diphenyl-2,2,3,3-tetramethylcyclotetrasiloxane by 29Si-NMR [120].

Figure 11.27 29Si-NMR spectrum of a silphenylene-siloxane block copolymer (67 mol % DMS), observed at 59.6 MHz [29] Reproduced with permission from E.A. Williams, J.H. Wengrovius, V.M. van Valkenburg and J.F. Smith, Macromolecules, 1991, 24, 1445, Figure 1. Copyright 1991, American Chemical Society

441


11.6.1.4 31P-NMR Chemical degradation of microporous polyurethane elastomer was investigated by using 1 H-, 13C- and 31P-NMR [121]. The degradation products polyester-polyurethane by phosphoric acid esters, such as (CH3CH2O)3PO and (ClCH2CH2O)3PO, gave a major 31 P-NMR peak at 0.24 ppm and -0.98 ppm as a septet, respectively. These peaks were assigned to phosphorous atom in the oligomer product with terminal phosphate groups, formed by exchange reaction as shown in Figure 11.28:

Figure 11.28 Structure of terminal phosphate groups of polyester-polyurethane by degradation with phosphoric acid esters

11.6.1.5 119Sn-NMR Polybutadiene and SBR can be anionically synthesised with tributyltin lithium (TBTLi) initiator in hexane/THF [122]. The 119Sn-NMR chemical shift of trialkyltin compounds is sensitive to the structure of substituent. For example, Bu3Sn-CH2CH=CHCH3 shows a signal at -17.79 and -13.69 ppm from Me4Sn for trans and cis isomers, respectively. Low molecular weight polybutadiene initiated with TBTLi and terminated with tributyltin chloride (TBTCl) exhibited signals reflecting the structure of butadiene units at both terminals as shown in Figure 11.29. It is clear that the signals a and e are observed only after termination with TBTCl and that the signal d is predominantly formed by initiation. Each signal was assigned to the sequences of 1,4 and 1,2 units, based on information from model compounds, as tabulated in Table 11.6. This indicates that the initiation starts from trans-1,4 and both 2,1 and 1,2 additions of butadiene units to tin-lithium bond and subsequent butadiene units can be added to form either cis-1,4, 2,1, 1,2 or trans-1,4 units. It is remarkable that no styrene unit was observed at the initiating terminal in the case of copolymerisation with butadiene and styrene.

442


Figure 11.29 119Sn-NMR of polybutadiene initiated with tributyltin Li (a) and terminated with tributyltin chloride (b), both 15% solutions in CDCl3, observed at 111.9 MHz [122] Reproduced with permission from W.L. Hergenrother, J.M. Doshak, D.R. Brumbaugh, T.W. Bethea and J. Oziomek, Journal of Polymer Science, Polymer Chemistry Edition, 1995, 33, 143, Figure 3. Copyright 1995, John Wiley & Sons

11.6.2 NMR of Swollen State and Latex

11.6.2.1 Swollen State NMR High-resolution NMR in solution requires the sample to be soluble in a solvent such that the various nuclear spin interactions can be averaged or removed by molecular microBrownian motions. Unfortunately, elastomers used in various applications are normally crosslinked materials and therefore not soluble in any solvent. Thus, solid state NMR with magic angle-spinning technique has been used with great success in the study of cured elastomers. However, this technique demands extended instrument facilities and expertise. It is well known that if a cured rubber sample were to be analysed by a conventional solution NMR spectrometer, very broad signals would be obtained. The resolution of the spectrum can be improved through swelling the rubber sample in a solvent. This

443

Spectroscopy of Rubbers and Rubbery Materials Table 11.6 Assignment of 119Sn-NMR signals of tributyltin-end-capped polybutadiene [122] Compound

ppm from Me4Sn

Peak Region

P—CH2CH=CHCH2 —(trans)

-18.17 to -18.34

f

-17.39 to -17.47

e

-15.97 to -16.33

d

CH=CH2 P—CH2CHCH2CHCH2CH CH=CH2

CH=CH2

CH=CH2 P—CHCH2— CH=CH2 P—CH2CHCH2 CH—

-15.46 to -15.51

c

-15.17 to -15.26

b

-14.60 to -14.76

a

CH=CH2 P—CH2CH— CH=CH2 P—CH2CH=CHCH2 —(cis)

Reproduced with permission from W.L. Hergenrother, J.M. Doshak, D.R. Brumbaugh, T.W. Bethea and J. Oziomek, Journal of Polymer Science: Polymer Chemistry Edition, 1995, 33, 143, Table 5. Copyright 1995, John Wiley & Sons

technique has been applied to the identification, analysis of molecular structure [123], and analysis of crosslink density of blends of vulcanised rubber [124]. Figure 11.30 shows the 13C-NMR spectra of polychloroprene rubber, obtained in solution (uncured), solid (cured), and swollen (cured) states. The rubber in swollen state has enhanced molecular mobility and therefore it showed improved spectral resolution. The enhanced spectral resolution makes it possible to calculate the ‘head-to-tail’ and ‘tail-to-head’ ratios for the triads composing of 1,4-trans unit in the polychloroprene rubber [123]. In the case of crosslinked EPDM, the spectral resolution was sufficiently good to carry out the monomer sequence analysis [123]. A similar technique was also applied to NBR. The triad sequence distribution of monomers was determined, by simplifying the copolymer to be consisting of acrylonitrile and butadiene, without discriminating between cis, trans, and vinyl isomeric units [123]. The results indicated the anticipated distribution in the rubber, reflecting the reliability of this technique in the analysis. Based on the

444


Figure 11.30 13C-NMR spectra of polychloroprene: (a) solution state, (b) solid state, (c) swollen state, observed at 100 MHz [123] Reproduced with permission from I. Fonao, L. Gonzalez, M.L. Jimeno and A. Marcos, Kautshuk und Gummi Kunststoffe, 1993, 46, 431, Figure 3. Copyright 1993, Hüthig GmbH

445

Spectroscopy of Rubbers and Rubbery Materials intensity ratios, the acrylonitrile content for several NBR rubbers was calculated and found to be close to the data supplied by the manufacturer. Analysis on NBR/PVC blends revealed that the weight fraction of each polymer was comparable to the known value [123]. For cured blends, however, the use of this technique to the composition analysis was less satisfactory due to the uneven distribution of crosslink density in the two rubber components [125]. Thus, the rubber with higher crosslink density gave smaller 13C-NMR signal area, resulting in less detection of this rubber in the blend. The signal broadening in 1H-NMR and 13C-NMR spectra, which caused by a reduction in molecular mobility in a rubber upon crosslinking, has been applied to the estimation of crosslink density in rubber blends [123]. The signal broadening effect could be correlated with crosslink density of the individual rubber, estimated by other methods such as equilibrium swelling or rheometer torque values [125]. For rubber blends prepared from unsaturated rubber, 1H-NMR can be used because the signal of the olefinic protons could be clearly observed without interference. In the case of blends containing saturated rubber, 13C-NMR will be the better choice because there are too many 1H-NMR signals in the alkyl proton regions [126,127]. The factors affecting the accuracy of this technique were also investigated [128].

11.6.2.2 Latex State NMR Many commercial elastomers are being produced in the form of latex. Characterisation of these samples in solution state might be difficult because these elastomers normally contain both sol and gel fractions. Therefore, attempts have been made to characterise directly these latices using high-resolution NMR spectroscopy. Early work on latex NMR involved the studies of acrylic latex [129,130]. The 13C-NMR signal is related to the polymer Tg and the temperature of measurement. The structure of polymer in the latex particles was also investigated in terms of temperature dependence of 13C-NMR linewidths [131]. The plot of line width against temperature yields a curve with two linear regressions. The intersection of the two linear plots is defined as breakpoint temperature. This temperature varies systematically with the interpenetrating network structure. The changes in 13C-NMR line-width are sensitive to the variation of crosslink density in an interpenetrating polymer network [131]. High mobility hydrophilic interfacial polymer chain will expand in the aqueous phase and therefore exhibit narrow lines in the NMR spectrum. The N-methylol acrylamide (NMA) in the functionalised butyl acrylate latex showed narrower carbonyl line than that of butyl acrylate [132]. The concentration of these hydrophilic groups is affected by the monomer concentration of NMA and also the timing of the addition of the monomer during the reaction.

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Recently, a novel method for determining the microstructure of crosslinked polybutadiene in latex using solution 13C-NMR technique was reported [133]. The surfactant and polymer concentrations in the latex were adjusted to give a good signal resolution of the latex sample, as indicated by half-width of the resonance peak at 32.7 ppm. Under these conditions, the S/N ratio was almost identical to that of sample in solution, as shown in Figure 11.31. The microstructure of sol and gel fractions in a radical initiated polybutadiene, determined by this technique, was similar to that of solution measurements.

Figure 11.31 13C-NMR spectra of polybutadiene containing 80% gel: (a) Solution measurement (10 w/v% CDCl3 solution), (b) latex measurement (10 w/v % dry rubber content), and (c) solid state measurement [133] Reproduced with permission from S. Kawahara, S. Bushimata, T. Sugiyama, C. Hashimoto and Y. Yanaka, Rubber Chemistry and Technology, 1999, 72, 2, 848, Figure 2. Copyright 1999, American Chemical Society

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Spectroscopy of Rubbers and Rubbery Materials The effect of particle size and spinning of the NMR tube were studied for the latex state 13 C-NMR of natural rubber latex fractionated by particle size [134]. High-resolution spectrum was obtained by measurement without sample spinning. The diffusion constant of Brownian motion was found to be a dominant factor governing the intensity and halfwidth of the signals. As the particle size decreased and temperature of measurement was raised, the intensity of signals increased and was comparable to the theoretical value, which was observed by the addition of triethylene glycol as an internal standard.

11.7 Conclusion High-resolution NMR has expanded its applicability to high molecular weight elastomers, with increasing the magnetic field from 60-100 MHz in the early stage to 500-750 MHz in modern 1H-NMR spectrometer. Improvement of the signal-to-noise ratio as well as the resolution of NMR spectrum makes it possible to detect very small amounts of sample and small signals derived from irregular structures of elastomers. This induced new application fields such as on-line HPLC-NMR measurement and analysis of end groups and branch points in addition to the conventional analysis of chemical composition, sequence distribution, microstructure of diene polymers. Applicability of many NMR pulse techniques, which have been widely applied in the case of low molecular weight compounds, to the assignment of small signals in elastomers increased the number of works in the field of structural characterization of elastomers. Now it is possible to analyse the structure of terminal groups, branch points, and crosslink points of elastomers with molecular weight higher than 105 by the use of modern NMR pulse techniques. These analyses provide direct information between the structure and physical properties as well as polymerisation mechanism including biosynthesis mechanism of natural rubber. It should be noted, however, the reliability of NMR analysis depends on the accuracy and precision of measurement in addition to the validity of signal assignment. Various traditional techniques, such as use of model compounds, partially deuterated polymers, 13 C-enriched polymers, and shift reagents, are very useful for signal assignment by combination with NMR pulse sequence techniques. Usage of nuclei other than 1H, 13C, and 19F will develop new application field of NMR of elastomers. Swollen state and latex state NMR will be just the method specialized to elastomers, because of no disturbance effect by the presence of gel fraction for these methods.

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12

Xe NMR of Elastomers in Blends and Composites

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Xe NMR of Elastomers in Blends and Composites Wiebren S. Veeman

12.1 Introduction to 129Xe NMR Spectroscopy of Materials

12.1.1 129Xe NMR Spectroscopy Most non-crystalline materials, and especially elastomers, readily absorb xenon (Xe) atoms. The Xe atoms can then be used to probe the local surrounding of the Xe atom in the material via nuclear magnetic resonance (NMR) spectroscopy. The direct information from standard NMR experiments usually comes from the chemical shift of the nucleus involved. The chemical shift of a nucleus is determined by the local surrounding of the atom (radius < approximately 1 nm). The additional information Xe NMR experiments can provide is because of the fact that, especially in elastomers, the Xe atom is very mobile. Depending on the type of NMR experiment, the characteristic time scale of a NMR experiment may be between 1 millisecond and 1 second and during this time a Xe atom may diffuse over several μm. An NMR investigation of Xe atoms in a material therefore provides information about the structure of the material averaged over a length scale of this magnitude. Compared to other branches of NMR, Xe NMR is unique in this sense. Several reviews on xenon NMR spectroscopy have appeared [1, 2], therefore in this introduction only the most important aspects will be discussed. The xenon atom has two isotopes which are suited for NMR studies, the 129Xe and 131Xe isotopes. For most studies 129 Xe is more convenient than 131Xe, while the former nucleus with spin I=1/2 does not have an electrical quadrupole moment. In some cases, however, the quadrupolar interaction can provide additional (spectral and relaxation) information. Here we will only consider the 129Xe isotope. For the probing of void spaces in and surfaces of materials Xe gas is ideal while Xe is an inert atom with a large (0.44 nm) and highly polarisable electron cloud. This is reflected by a very wide 129Xe chemical shift range (over 7500 ppm), although the larger part of this range results from Xe covalently bonded in compounds. The effective chemical shift

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Spectroscopy of Rubbers and Rubbery Materials range for Xe atoms sorbed in materials is about 300 ppm. This is of course still very large in comparison to, for instance, the 1H chemical shift range. In the experiments described here, the material investigated is contained in a glass tube in a Xe atmosphere at ca. 1,000,000 Pa. As a result each Xe spectrum usually contains at least two lines, one from Xe in the gas surrounding the material but inside the NMR coil, and one for Xe, either absorbed in the material or adsorbed at pore and outer surfaces of the material (Figure 12.1). Since the Xe NMR resonance of the free gas is used as the spectral reference, the NMR of Xe gas will be discussed shortly.

12.1.2 NMR of 129Xe in the Gas Phase Because of the collisions between the Xe atoms in the gas phase, the chemical shift of 129 Xe in the gas phase is temperature and pressure dependent. Quantitatively, a relationship between the 129Xe chemical shift in ppm and the density of the gas has been found to exist [3]:

δ ( T , p ) = δ 0 + δ1( T )ρ + δ 2 ( T )ρ 2 + ...........

(12.1)

where δ0 is a reference shift set to 0 ppm; δ1, δ2,..... are coefficients and ρ is the density. When the density is expressed in the unit Amagat (1 Amagat is the density of the gas at 1 atmosphere at 0 °C and corresponds to approximately 2.7 x 1019 atoms/cm3, assuming an ideal gas) the coefficients at 25 °C are δ1 = -0.548 ppm/amagat and δ2 = -0.169 x 10-3 ppm/amagat2. For the pressures used in the experiments described here higher terms in Equation 12.1 need not be considered.

Figure 12.1 A typical 129Xe NMR spectrum of a polymer at a temperature above the glass transition temperature (Tg) (here ethylene-propylene diene terpolymer (EPDM)) in a ca. 1,000,000 Pa Xe atmosphere. The signal of the free gas is used as an internal chemical shift resonance

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For our experiments the sample is located in a sealed glass tube in a Xe atmosphere and the signal of the free Xe gas is taken as the 0 ppm reference. The initial pressure is carefully adjusted to 1,000,000 Pa, but slight deviations (50,000 Pa) may occur during the sealing process. Using Equation 12.1 and assuming that the Xe gas behaves like an ideal gas, a pressure variation of 50,000 Pa leads to chemical shift variations of ±0.2 ppm. In view of the width of the lines, this variation of the chemical shift reference with pressure is neglected here. Materials at temperatures above the glass transition temperature (Tg) absorb Xe readily, therefore the fact that after sealing of the sample tube the density of the Xe gas could decrease with time when more Xe is absorbed by the material is neglected. For samples which absorb Xe slowly, like some polymers below their Tg, this effect may not be negligible. Also for temperature dependent measurements the variation of the coefficients in Equation 12.1 with temperature should be taken into account [4, 5, 6], but in view of the widths of the resonance lines investigated here as a function of temperature, the temperature dependence of the reference has also been neglected.

12.1.3 129Xe NMR of Polymers Before describing 129Xe NMR experiments on polymers and polymer composites, it is worthwhile to show with a few selected examples, some of the general aspects of 129Xe NMR of polymers. The first question that needs to be addressed is where sorbed Xe atoms are located in a polymer material, or rather where Xe atoms are not located. With a diameter of 0.44 nm the Xe atom is clearly larger than the interchain distance for most crystalline polymers. In general it means, that if Xe atoms are found in crystalline domains, they must reside in areas which contain defects. In the examples studied, no evidence of Xe in crystalline domains of polymers was ever found. A clear example for the absence of Xe in crystalline polymers is shown in Figure 12.2 for highly stretched polyethylene (PE) fibres. From X-ray experiments it is known that the crystallinity of the PE fibres increases with stretching. In Figure 12.2 the NMR signal of 129Xe absorbed in the PE disappears with increasing stretching, which shows that the Xe is practically absorbed only in the amorphous domains, and possibly in the interface between crystalline and amorphous domains, of PE. The absence of Xe in crystalline polymers is also due to the fact that the polymer chains in crystalline polymers are quite rigid. The energy needed to deform the chains enough so that a Xe atom can be incorporated between the chains, is too high.

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Figure 12.2 129Xe NMR spectra of PE fibres as a function of stretching. The Young’s modulus increased from 38 GPa for PE1 to 131 GPa for PE5

The mobility of the chains plays an important role for the line width of the Xe resonance in amorphous polymers. Figure 12.3 shows the line width of the 129Xe resonance of 129 Xe absorbed in polymethylmethacrylate (PMMA) as a function of temperature [7]. The line width drastically reduces when the Tg is approached. The mobility of the polymer chains in amorphous domains affects the 129Xe line width while it makes it possible for the Xe atom to move fast from one location to another, thereby averaging out local differences in Xe chemical shift and dipolar interactions with proton spins [8]. The chain mobility also influences the rate of Xe absorption in an amorphous polymer, because it was noted that for PMMA at room temperature it takes several weeks before enough Xe is absorbed in the material so that a NMR signal can be detected [7]. From Figure 12.3, it can be concluded that room temperature 129Xe NMR spectroscopy is especially well suited for the investigation of amorphous domains in polymer materials

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Figure 12.3 The line width of the 129Xe NMR resonance of PMMA as a function of temperature. The sharp decrease of the line width for temperatures below Tg shows that the mobility of the Xe atoms already increases at temperature for below the Tg Reproduced with permission from K. Sperling-Ischinsky and W.S. Veeman, Journal of the Brazilian Chemical Society, 1999, 10, 293, Figure 11. Copyright 1999, The Brazilian Chemical Society

for which the Tg is clearly below room temperature. In such materials the Xe atoms are very mobile, therefore the Xe diffusion coefficient can also be considered as a characteristic parameter of the material. The determination of the Xe self-diffusion coefficients allows the calculation of the root mean square displacement of diffusing xenon atoms with time in the material.

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12.1.4 129Xe Pulsed Field Gradient Echo (PFGE) Spectroscopy The PFGE experiment is a NMR imaging experiment [9, 10], in which pulsed magnetic field gradients are used to make the Larmor precession frequency of the 129Xe spins during the time of the gradient pulses dependent on the spatial coordinates parallel to the direction of the field gradient, The scheme of the radiofrequency (rf) and magnetic field pulses is shown in Figure 12.4. In addition to the rf pulses (representing a stimulated spin echo experiment) two pulsed gradients, identical in amplitude g, width δ and direction, and separated by a time Δ >> δ are applied. The first gradient pulse causes a phase jump of the magnetisation transverse to the static external magnetic field. This phase jump linearly depends on the coordinate of the spin position in the direction, e.g., x, of the field gradient. For spins which do not diffuse during the time change (Δ) between the two gradient pulses, or which at the time of the second gradient pulse have returned to the same x-coordinate value they had at the time of the first pulse, this phase jump is exactly compensated by the second (identical) gradient pulse thanks to the spin echo sequence. They do not suffer an overall phase loss and contribute fully to the total intensity of the spin echo signal. For spins which did diffuse from their original position (or more accurately who changed their x-coordinate value) during the time change, the phase jump their magnetisation vector acquired during the first pulse is not compensated by the phase jump during the second pulse and their contribution to the echo intensity is decreased.

Figure 12.4 The pulse scheme of the three-pulse echo sequence to determine Xe diffusion coefficients in polymers and other porous systems. The shaded areas are magnetic field gradient pulses with amplitude g and length δ. The time between the two gradient pulses Δ determines the time during which the diffusion path length is determined (the diffusion time) Reproduced with permission from F. Junker and W.S. Veeman, Macromolecules, 1998, 20, 7010, Figure 1. Copyright 1998, American Chemical Society

462

129


In the limit that δ C*), or in the molten state when the molar mass is greater than the critical mass for entanglements, there is often observed two separate decays with largely-different time constants of decay: M x, y (t) M x, y (0)

502

= f1R1(t) +f2 R 2 (t)

Swollen Rubbery Materials: Chemistry and Physical Properties Studied by NMR Techniques The fractions of protons decaying according to relaxation functions R1 and R2 are given by f1 and f2. In molten polymers this relationship has long been exploited to provide a measure of the crosslink density in many polymer systems [64-83]. The form of the decay functions has been the subject of much discussion, however, it is often observed that R1 and R2 can be approximated by simple exponential decay functions. It is generally accepted that the protons with short relaxation times are those directly attached to or adjacent to crosslink points. As an example Figure 13.4 shows the decays of transverse

Figure 13.4 1H spin-spin relaxation decays recorded using the Hahn-echo sequence for polyisoprene at 150 °C, crosslinked with radiation doses of 0 (●), 0.3 (O), 0.58 (▲), 1.5 (■), 2.25 (❑) and 5.2 kGy [65] Reprinted with permission from R. Folland and A. Charlesby, Polymer, 1979, 20, 211, Figure 1, Copyright 1979, Elsevier Science

503

Spectroscopy of Rubbers and Rubbery Materials magnetisation in polyisoprene crosslinked with increasing dose of gamma radiation [65]. Clearly the rates of decay increase with increasing radiation dose. In addition the proportion of protons decaying more slowly decreases with dose. When the irradiation dose is corrected for a ‘virtual dose’ to account for the presence of physical entanglements present prior to irradiation, the fraction of slowly decaying protons agrees very closely with the soluble fraction measured by Soxhlet extraction. Although this and similar studies were performed on molten materials, the same information could be obtained from swollen samples. Indeed the range of crosslink densities that could be examined using this method would be extended by swelling in a deuterated solvent, since the difference between the rates of fast and slow relaxation would be increased [84], and hence resolution of the two decays would become more facile. The effect of swelling on the rates of transverse relaxation has been examined by Fukumori and co-workers [84]. T2 relaxation times were measured for carbon black-filled nitrile rubber (NBR) during swelling with carbon tetrachloride. The relaxation decay was resolved into a two exponential process, the slower of which became more predominant during swelling. The fraction of protons relaxing slowly, i.e., the more mobile chains, was found to increase with increasing swelling ratio, and was therefore directly proportional to the squared root of swelling time, as predicted by Fickian models of diffusion. In addition there is evidence of more effective rubber-filler interactions, and restricted pathways to diffusion with increased filler content persistent at long diffusion times. While it is generally accepted that the slow relaxation decay in Equation 13.2 can be represented by an exponential decay, it is well known that the rapid decay of magnetisation conforms more closely to a gaussian decay function. Folland and Charlesby [65] recognised this but were more concerned with changes in the relative proportions of the rapidly and slowly decaying functions. In 1976 Gotlib and co-workers [85] developed an expression for the second moment of the NMR line shape by considering the anisotropic motion of constrained polymer chains. They were the first to recognise that the average moment will not depend on the details of the chain motion, but rather the average conformations of the polymer chains. In their case the conformational mean was calculated using the results of Kuhn and Grün for freely jointed chains [86]. They were able to show that the plateau value of T2 of crosslinked polystyrene swollen in toluene was inversely proportional to the crosslink density. Fry and Lind [76] later showed that these calculations could be applied to heavily crosslinked epoxy resins. More recently, Litvinov and coworkers have used these calculations to determine the density of chemical crosslinks and chain entanglements in cured EPDM [87], the interactions between carbon black and polymer chains in filled EPDM [88], and the average length of amorphous chains spanning inter-lamellar spaces in semi-crystalline polyvinyl chloride [89].

504

Swollen Rubbery Materials: Chemistry and Physical Properties Studied by NMR Techniques Simon and co-workers [71-75] later developed a variation of this model of T2 relaxation in which the faster decay process arises from an ‘inter-crosslink’ component of the network, in which motion is rapid and anisotropic and does not eliminate completely the dipolar interactions. The rapid motion, described by a correlation time τf, occurs simultaneously with a slower isotropic motion of the complete inter-crosslink segment described by the correlation time τs. The longer relaxation decay arises from dangling chain ends and is exponential in nature. The total decay of transverse magnetisation if then given by:

M x, y (t) M x, y (0)

= A.exp[

−t −t −t t − qM 2 τ s2 (exp( ) + − 1)] + B.exp[ ] τs τs T2 T2

where M2 is the second moment of the rigid-lattice line shape, q is the fraction of the rigid-lattice second moment unaveraged by the anisotropic motion, and 1/T2 = M2τf. An additional long relaxation decay having approximately 5% of the total signal amplitude was ascribed to uncrosslinked or sol chains. It was argued that the parameter q, the degree of averaging of the rigid-lattice second moment, could be related to the molecular weight between crosslinks Mc as:

Mc =

3cMru 5n q'

where q′ is parameter q for a crosslinked network minus the contribution from physical entanglements (qo) measured for an uncrosslinked material. This method was recently applied by Menge and co-workers [90] to the study of swollen PDMS networks. While the analysis of the transverse relation behaviour of the unswollen network conformed well with the theory of Simon and co-workers, on swelling more complex behaviour was observed. For weakly swollen materials the parameter q, the extent of averaging of the second moment decreased, however, for strongly swollen networks relaxation became more efficient at short times. This was indicative of increased anisotropy of motion of inter-crosslink segments due to deformation of the network. The relationship between Mc and q thus breaks down in this situation. Application of the empirical relationship proposed by Schmidt and Cohen-Addad [91] demonstrates the departure at higher swelling ratios from Gaussian behaviour. In more recent years extensive work by Cohen-Addad [92-95] and Brereton [96-98] has lead to a more comprehensive understanding of the effect of chain entanglements on the decay of transverse relaxation in the NMR experiment. In their scale-invariant model, the polymer chain is considered to consist of a series of sub-units, the smallest of which

505

Spectroscopy of Rubbers and Rubbery Materials consists of a small number of monomer units. The dipolar interactions experienced by the protons are scaled by the number of units in these elemental units. It was further shown [96] that in the limit of fast local motion, the dipolar interactions can be further scaled over larger segments which consist of a number of the smaller units, and in crosslinked polymers these units are connected by crosslinks and chain entanglements. It was found that the decay of transverse magnetisation is insensitive to the correlation time of motion of these larger units, but depends on the number of Rouse units between the crosslink points, and hence the molar mass between crosslinks. These methods have been applied with success to molten linear polymers, however, the effects of free chain ends in crosslinked polydisperse materials limits its applicability [99]. This topic is discussed in detail in Chapter 8.

13.5.2 Estimation of Crosslink Density from NMR Linewidths

13.5.2.1 1H NMR Linewidths In 1989 Loadman and Tinker [100] published a study of the variation of linewidth in the 1 H NMR spectra of NR and NBR with crosslink density. In this and subsequent papers their aim was to measure the crosslink density in blends of two rubbers. For this reason they developed an empirical method for determining crosslink density from the peak widths. The width was estimated from the ratio of the height of the peak at +/- 0.205 ppm from the centre of the peak to the maximum height of the peak. This parameter, which they called H%, was measured on the low field side of the olefinic peak of NBR, and on the high field side of the olefinic peak of NR, to minimise interference from the overlapping peaks in the blends. None-the-less, there is significant overlap in spectra of blends which must be corrected for by reference to the spectra of the pure crosslinked swollen rubbers. An iterative procedure is described in this initial paper [100] which results in a measure of the equilibrium swelling of each component. The relationship between H% and the volume fraction of rubber in swollen NR is shown in Figure 13.5. Although it is not possible, given the information provided, to construct a comprehensive theory relating the mechanism of line broadening with the extent of network formation, the value of this work lies in the ability to provide a measure of crosslink densities in blended and filled systems. In a second paper Brown and Tinker [101] examined the effects of a number of parameters, such as the accelerator used in the sulfur curing, peroxide versus sulfur curing, and swelling ratio at constant crosslink density. The results for cis-polyisoprene show that the value of H% is independent of accelerator. However, lower values of H% were seen for peroxide-cured materials. The results for BR were independent of curant; the reasons

506

Swollen Rubbery Materials: Chemistry and Physical Properties Studied by NMR Techniques

Figure 13.5 The relationship between the linewidth parameter H% and swelling ratio for swollen NR [100]. The measurements were made in deuterochloroform at 23 °C [100] Adapted with permission from M.J.R. Loadman and A.J. Tinker, Rubber Chemistry and Technology, 1989, 62, 234, Figure 4. Copyright 1989 Rubber Division, American Chemical Society

for this were not clearly explained. In another series of experiments, measurements were made on partially swollen BR rubbers; a weaker dependence of H% on the swelling ratio compared to the equilibrium swelling was observed. It was concluded that the method is more sensitive to the crosslink density rather than the degree of swelling and is hence a robust method for determining crosslink density. The presence of carbon black filler increases the linewidths in the spectra from both phases of a blend, and as a result comparison could only be made of the relative crosslink density [102]. Later the authors showed that H% in the presence of filler in blends of NR and EPDM was linearly related to H% in the unfilled blend [103]. In blends of NR and BR a small excess of the sulfur reacted in the NR phase compared with the BR. Measurement in unfilled blends allowed crosslink densities to be determined, and confirmed the greater yield of crosslinking in the NR, is shown in a plot of H% against crosslink density [104].

507

Spectroscopy of Rubbers and Rubbery Materials Separate work on a Fourier transform (FT) NMR spectrometer revealed the presence of two peaks from the solvents from both within and outside the swollen gel [104]. See below for a discussion of the origin of the NMR linewidths. The chemical shift of the olefinic peaks was found to shift down field with increasing crosslink density, and hence a modified method for determining H% was introduced. In this paper they also introduced the first 13C NMR measurements of swollen rubber blends, and again found a systematic increase in linewidth with increasing crosslink density. The higher resolution in the 13C spectrum compared with 1H NMR allows the potential of more detailed information on rubber mixtures. The effect of partitioning of curatives on the crosslinking reactions in NR/BR blends was explored in more detail in a later paper [105]. The ultimate extent of curing observed in the individual phases of the blends was identical to that obtained for the pure components, however, for the blend faster curing was initially observed in the BR phase. This was related to the greater affinity of sulfur and accelerator for BR compared with NR. Other systems examined include blends of epoxidised NR and cis-BR [106], NR blended with cis-BR [107] and NR blended with EPDM [108]. The work of Tinker and co-workers has been discussed at length by Cook [109].

13.5.2.2 13C NMR Linewidth Before leaving this subject it is of value to discuss the possible use of swollen-state 13C NMR to determine crosslink densities in swollen rubbers. As briefly mentioned previously, Brown and co-workers [104] showed that measurements of the crosslinking density could be obtained using 13C NMR. They examined in much greater detail the curing of blends of NR and EPDM, including chemically-modified EPDM, in a later paper [108]. The much greater resolution afforded by 13C spectroscopy has the potential to provide more detail of the mechanism of reactions occurring during curing. The work of Whittaker and O’Donnell [110] on radiation-crosslinked ethylene-propylene rubber (EPR) showed that the overall 13C NMR signal intensity decreased with increasing crosslink density (Figure 13.6). This was also observed earlier by Ford and co-workers for crosslinked polystyrene (PS) [111]. The scalar decoupling used in most solution-state NMR experiments is insufficient to remove strong dipolar couplings to 1H nuclei present at 13C nuclei involved directly in crosslinks or adjacent to crosslinks. As a consequence the NMR signal is observable only from nuclei well removed from the crosslink points. Measurement of the chemical structure of crosslinks by solution-state NMR is therefore restricted to materials crosslinked to levels below the gel point [110, 112-114]. It also follows that for higher crosslink densities the 13C NMR signal is not representative of the whole sample, but that rather the motion of the chains well removed from the crosslinks reflects the crosslink density of the entire sample.

508


Figure 13.6 Decrease in 13C NMR signal intensity (●) and soluble polymer fraction (■) for crosslinked EPR containing 36 mol% propylene, as a function of radiation dose [110]. The NMR experiments were conducted in deuterated chloroform at 25 °C. The solid lines are drawn as a guide to the eye Adapted with permission from J.H. O’Donnell and A.K. Whittaker, Polymer, 1992, 33, 62, Figure 3, Copyright 1992 with permission from Elsevier Science

13.5.2.3 Origin of Changes in Linewidth The origin of the linewidths in the 1H and 13C spectra of crosslinked polymer has received some attention. Mohanraj and Ford [111] found an approximately linear increase in 13C linewidth with magnetic field strength (2.3 T and 7 T) for swollen crosslinked PS, and therefore suggested the linewidth was determined by chemical shift dispersion, although at higher crosslink densities a smaller increase in linewidth was recorded. It was concluded that other mechanisms, for example dipole-dipole couplings could become important for highly-immobilised gels. This is certainly reasonable when using standard solutionstate NMR equipment. Bain and co-workers [115] studied the field dependence of 13C linewidths for polyacrylates crosslinked to low levels. The linewidth was directly proportional to field strength, consistent with a significant contribution from a dispersion of chemical shifts arising from new chain structures. In addition the authors used the Delays Alternating with 509

Spectroscopy of Rubbers and Rubbery Materials Nutations for Tailored Excitations (DANTE) pulse sequence [116] to selectively irradiate a small band of frequencies in the centre of the methylene resonance. Despite insufficient details of this experiment being provided this result was also consistent with the linewidth arising from inhomogeneous broadening. More recently, however, Stover and Frechet [117] have shown that MAS dramatically reduces the linewidths in 13C spectra of crosslinked PS gels. MAS is effective in averaging weak dipole-dipole couplings and broadening due to magnetic field inhomogeneities. In 1992 Brown and co-workers [104] showed that two peaks were observed in the 1H NMR spectra for the small molecules both inside and outside the swollen gel. This is indicative of a difference in bulk susceptibility and hence local field in these two environments. It is likely therefore that this mechanism also contributes to the NMR linewidths in these heterogeneous systems. Line broadening due to susceptibility effects would scale linearly with field, as observed by Errade and co-workers [118], would lead to an inhomogeneous line susceptible to hole burning [115], and be averaged by MAS [117]. At higher crosslink densities dipolar coupling would become increasingly important. To demonstrate that this issue is far from being resolved, a study of linear unswollen polyisoprene by English [119] concluded that the dominant mechanism of broadening in this case was dipolar in origin resulting from anisotropic motion of chain segments. Given the above discussion it is considered that these two effects, namely weak dipolar interactions and local field inhomogeneities, contribute most significantly to linewidths in the 1H and 13C spectra at moderate crosslink densities [120].

13.5.3 High-Resolution 13C MAS NMR of Rubbers It has recently been recognised that high resolution NMR spectra of molecules attached to swollen, crosslinked polymer beads can be obtained by the method of MAS. A number of instrument manufacturers are marketing MAS probes for both 1H and 13C observation designed for highly uniform static magnetic fields. The main potential area of application of this technology is the field of combinatorial chemistry, in which small molecules are synthesised on polymer matrices, and are required to be characterised before removal. It follows that high resolution MAS NMR of swollen polymers will become an important tool for characterisation of new structures and surface grafting in heterogeneous rubbers. As an example, Doskocilova and co-workers [121] have recently used swollen-state 1H and 13C MAS NMR to measure changes in structure of BR on thermal degradation. The method of cross-polarisation was used to obtain spectra identical to those measured using single pulse excitation. In highly heterogeneous materials cross-polarisation will tend to favour signals from more rigid regions (at shorter cross-polarisation contact times), while the single pulse experiment is biased towards 13C nuclei with shorter T1 relaxation times, that is in this case more mobile carbon nuclei. This result provided evidence that the material was homogeneous in nature.

510


13.6 Summary It is clear from the examples given in this review that there are many reasons for examining the structure of crosslinked rubbers by swollen-state NMR. The first of these is the increased resolution in the NMR spectrum compared with the molten state, as a result in the more effective averaging of the dipole-dipole couplings between nuclear spins. With conventional high-resolution solution NMR equipment it is possible to determine the crosslink structure and concentrations of crosslinks prior to gelation. The demands on spectrometer time for these experiments are prodigious since the concentrations of carbons or proton nuclei participating in crosslinks are by definition less than one per polymer chain. However, the new information obtained on the mechanisms of crosslinking fully justifies this investment in time. At higher crosslink concentrations two approaches can be adopted. Firstly the NMR experiment can be designed to measure a parameter sensitive to changes in the segmental motion of the bulk sample. Examples of these are measurements of 1H spin-spin relaxation times, or widths of lines in the spectra. Since the NMR signal does not arise from the entire sample under solution-state conditions it is necessary to construct a master curve relating the NMR parameter to the crosslink density. This method is especially useful for mixtures of rubbers or for filled rubbers. A second approach is to use the methods of MAS and high-power proton decoupling to increase the efficiency of averaging of the residual dipole-dipole couplings, and broadening due to magnetic field inhomogeneities, and hence obtain a high-resolution spectrum of the whole sample. This is expected to be an important method in future studies of gel structure. An additional advantage of the reduced NMR linewidths and rates of transverse relaxation is the possibility of higher sensitivity in NMR imaging experiments. This advantage is of course most evident in systems where the transverse relaxation times are reduced to very short values by high local crosslink densities or the presence of fillers. Examples have been given where images are obtained of either the plasticised rubber phase or of the solvent itself. The measurement of diffusion coefficients in the presence of a gradient in solvent concentration by NMR imaging is also an important experiment with practical significance, as is the measurement of self-diffusion coefficients by PFG NMR. An understanding of the effect of crosslinking on local chain motion in polymers has been developing over the past 20 years. The strong agreement between values of correlation times for segmental motion of uncrosslinked chains determined by NMR and other methods, and the clear relationship with free volume theory underlines the importance of the NMR experiments. It is evident that the presence of permanent crosslinks in polymer chains leads to a pronounced motional heterogeneity compared with linear chains in solution or in the bulk state, for which a sound theoretical understanding of the mechanism

511

Spectroscopy of Rubbers and Rubbery Materials of chain reorientation has previously been established. The full application of this knowledge to crosslinked rubbers has yet to be achieved, and promises to shed light on the effect of crosslinking on segmental chain motion.

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Swollen Rubbery Materials: Chemistry and Physical Properties Studied by NMR Techniques 87. V.M. Litvinov, W. Barendswaard and M. Van Duin, Rubber Chemistry and Technology, 1998, 71, 105. 88. V.M. Litvinov and P.A.M. Steeman, Macromolecules, 1999, 32, 8476. 89. W. Barendswaard, V.M. Litvinov, F. Souren, R.L. Scherrenberg, C. Gondard and C. Colemonts, Macromolecules, 1999, 32, 167. 90. H. Menge, S. Hotopf, S. Ponitzsch, S. Richter, K.F. Arndt, H. Schneider and U. Heuert, Polymer, 1999, 40, 5303. 91. C. Schmit and J.P. Cohen-Addad, Macromolecules, 1989, 22, 142. 92. J.P. Cohen-Addad, Journal of Physics, 1982, 43, 1509. 93. J.P. Cohen-Addad and R. Dupeyre, Polymer, 1983, 24, 400. 94. J.P. Cohen-Addad, Polymer, 1983, 24, 1128. 95. J.P. Cohen-Addad and J. Guillermo, Journal of Polymer Science: Polymer Physics Edition, 1984, 22, 931. 96. M.G. Brereton, Macromolecules, 1990, 23, 1119. 97. M.G. Brereton, Journal of Chemical Physics, 1991, 94, 2136. 98. M.G. Brereton, Macromolecules, 1991, 24, 2068. 99. M.E. Reis, M.G. Brereton, P.G. Klein and P. Dounis, Polymer Gels and Networks, 1997, 5, 285. 100. M.J.R. Loadman and A.J. Tinker, Rubber Chemistry and Technology, 1989, 62, 234. 101. P.S. Brown and A.J. Tinker, Journal of Natural Rubber Research, 1990, 5, 286. 102. P.S. Brown and A.J. Tinker, Kautschuk und Gummi Kunststoffe, 1995, 48, 606. 103. M.D. Ellul, J. Patel and A.J. Tinker, Rubber Chemistry and Technology, 1995, 68, 573. 104. P.S. Brown, M.J.R. Loadman and A.J. Tinker, Rubber Chemistry and Technology, 1992, 65, 744.

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Spectroscopy of Rubbers and Rubbery Materials 105. S.A. Groves and A.J. Tinker, Journal of Natural Rubber Research, 1996, 11, 125. 106. P.S. Brown and A.J. Tinker, Journal of Natural Rubber Research, 1991, 6, 87. 107. P.S. Brown and A.J. Tinker, Journal of Natural Rubber Research, 1993, 8, 1. 108. P.S. Brown and A.J. Tinker, Journal of Natural Rubber Research, 1996, 11, 227. 109. S. Cook, Kautschuk und Gummi Kunststoffe, 1999, 52, 350. 110. J.H. O’Donnell and A.K. Whittaker, Polymer, 1992, 33, 62. 111. S. Mohanraj and W.T. Ford, Macromolecules, 1985, 18, 351. 112. J.C. Randall, F.J.Zoepfl and J. Silverman, Makromolekulare Chemie, Rapid Communications, 1983, 4, 149. 113. F. Horii, Q. Zhu, R. Kitamaru and H. Yamaoka, Macromolecules, 1990, 23, 977. 114. Q. Zhu, F. Horii, R. Kitamaru and H. Yamaoka, Journal of Polymer Science: Polymer Chemistry Edition, 1990, 28, 2741. 115. A.D. Bain, D.R. Eaton, A.E. Hamielec, M. Mlekuz and B.G. Sayer, Macromolecules, 1989, 22, 3561. 116. G.A. Morris and R. Freeman, Journal of Magnetic Resonance, 1978, 29, 433. 117. H.D.H. Stover and J.M.J. Frechet, Macromolecules, 1991, 24, 883. 118. L.A. Errede, R.A. Newmark and J.R. Hill, Macromolecules, 1986, 19, 651. 119. A.D. English, Macromolecules, 1985, 18, 178. 120. D. Doskocilova and B. Schneider, Pure and Applied Chemistry, 1982, 54, 575. 121. D. Doskocilova, J. Straka and B. Schneider, Polymer, 1993, 34, 437.

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14

Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials Dan E. Demco, Siegfried Hafner and Hans W. Spiess

14.1 Introduction An important objective in materials science is the establishment of relationships between the microscopic structure or molecular dynamics and the resulting macroscopic properties. Once established, this knowledge then allows the design of improved materials. Thus, the availability of powerful analytical tools such as nuclear magnetic resonance (NMR) spectroscopy [1-6] is one of the key issues in polymer science. Its unique chemical selectivity and high flexibility allows one to study structure, chain conformation and molecular dynamics in much detail and depth. NMR in its different variants provides information from the molecular to the macroscopic length scale and on molecular motions from the 1 Hz to 1010 Hz. It can be applied to crystalline as well as to amorphous samples which is of particular importance for the study of polymers. Moreover, NMR can be conveniently applied to polymers since they contain predominantly nuclei that are NMR sensitive such as 1H and 13C. While well-established for liquid-like samples such as macromolecules in solutions [7], the applications of NMR to solid or solid-like polymers is more demanding because of the presence of anisotropic interactions that complicate the analysis of the results. Several techniques for the removal of these interactions have thus been developed and are nowadays in a state where they can be routinely applied. Non-averaged anisotropic interactions on the other hand provide valuable information that is lost in the solution state. Thus, while it is often necessary to remove the anisotropic interactions, in many cases one would simultaneously like to preserve them in order to exploit their information content. This is where two-dimensional (2D) spectroscopy comes into play, for example, by correlating one dimension where the anisotropic interactions are preserved with a (highresolution) dimension where they are removed. That is, both objectives can be achieved within one two-dimensional experiment. Like conventional one-dimensional (1D) NMR techniques, 2D techniques were first applied to liquid or solutions, where they provided invaluable information for the structure assignments in biomolecules [7]. Later, after the

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Spectroscopy of Rubbers and Rubbery Materials experimental difficulties related to the NMR of solids have been overcome, they were also applied to solid polymers and elastomers [1-6]. From the viewpoint of NMR, elastomers and other viscoelastic polymers above their glass-transition temperature (Tg) exhibit both, solid-like and liquid-like features. Whereas the segmental motions give rise to the liquid-like behaviour, the presence of permanent or non-permanent crosslinks leads to residual dipolar couplings, that are responsible for the solid-like properties [8]. While this promises that both properties can be exploited, the application of 2D techniques to viscoelastic materials has to deal with the difficulties related to both, rigid and mobile samples. Compared with the wealth of applications in solution and in solid state, it has not been widely applied to viscoelastic polymers, although it can provide information on such important fields as the chain dynamics in elastomers, the local structure, residual couplings (induced by chemical crosslinks and topological constraints), dynamic order parameters, inter-nuclear distances, intermolecular interactions (which are important, for example, for the miscibility), the effects of fillers on molecular motions, segmental orientation under mechanical stress and others. After a brief introduction of the basic tools of NMR in Section 14.2, the 2D techniques that have been already applied to rubbers or viscoelastic materials (Section 14.3) will be reviewed. After briefly introducing each of the techniques, a more detailed overview of the applications and a discussion of some of the highlights will be given. This structure, where all information on a given method is presented within one section enables the interested reader to decide more easily which of the techniques might be most useful to them. NMR imaging which can be considered as a special form of 2D NMR will not be discussed but the interested reader is referred to the corresponding chapter in this book.

14.2 Basics of NMR in Viscoelastic Polymers

14.2.1 Anisotropic Spin Interactions The information that can be extracted from solid-state NMR spectra is encoded via spin interactions such as the chemical shielding, the quadrupolar interaction and the homoand hetero-nuclear dipolar interactions [1,9-10]. Some knowledge of the spin interactions that determine the features of the spectra are thus of prime importance. A common characteristic of the relevant spin interactions is that they are anisotropic and can be described by second-rank tensors. The resulting orientation-dependent NMR frequency is of the following form [1,9]:

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Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials

ω(α, β) = ω L +

δ (3 cos2 β − 1 − η sin2 β cos 2α) 2

(14.1)

where ωL is the Larmor frequency at the isotropic chemical shift value and the other terms reflect the deviations due to angular-dependent contributions. δ represents the strength of the anisotropic interactions, η (0 ≤ η ≤ 1) is the asymmetry parameter which describes the deviation from axial symmetry and α, β are the polar angles that relate the orientation of the principle-axes system of the interaction tensor with the main magnetic field. The most important interaction for chemical information is the chemical shift. It results from the shielding of the magnetic field at the position of the nucleus by the electrons. The Larmor frequency is thus shifted by an isotropic contribution ωL and by angular dependent contributions (polar angles α and β). Assuming an equal probability for all directions one can calculate the powder average and a broad powder spectrum is obtained which reflects the chemical-shift anisotropy. For the case of an axially symmetric chemicalshift tensor, η is zero and the angular-dependent term in Equation 14.1 simplifies to δ (3cos2β-1)/2. Usually the chemical shift spans about 10 ppm for 1H and 200 ppm for 13 C nuclei. For abundant nuclei with spin 1/2, the spectrum is often dominated by heteronuclear or homonuclear dipolar interactions, i.e., the interactions between the magnetic moments of two neighbouring spins. In this case there is no isotropic contribution and η is zero, so that Equation 14.1 simplifies correspondingly. For a two-spin system one obtains a spin Hamiltonian of the form:

rr ⎛ μ ⎞ γ γ h ⎛ 3 cos2 θ − 1⎞ (3I1z I2z − I1 I2 ) H D = ⎜ 0 ⎟ 1 32 ⎜ ⎟ ⎝ 4 π ⎠ r12 ⎝ 2 ⎠

(14.2)

where r12 is the magnitude of the r vector connecting the two spins, θ is the angle of this vector to the magnetic field, the Ii are spin operators and the γi (i=1,2) are the magnetogyric ratios of the spins. That is, the strength of the resulting line splitting depends strongly on the distance between the two spins, so that information about distance can be extracted from such spectra. Homonuclear interaction (equivalent spins with γ1 = γ2 = γ) and heteronuclear interaction (non-equivalent spins with γ1≠γ2) have r r to be distinguished. In the latter case, the flip-flop term which is part of the product I1I2 in Equation 14.2 can be neglected. For a powder sample one has again to take into account all angles β and thus obtains the so-called Pake spectrum [1,9] with a considerable anisotropic linebroadening of up to 50 kHz for homonuclear and up to 25 kHz for heteronuclear dipolar interaction. Since the dipolar coupling is a through-space interaction, however, in principle

521

Spectroscopy of Rubbers and Rubbery Materials the sum has to be evaluated over all possible pair interactions. This, and the presence of molecular motion, leads to considerable complications and is responsible for the experimental finding that in practice not a Pake spectrum but a relatively structureless line-shape is obtained. In case of a deuterated sample (spin 1 case), the spectra are usually dominated by the quadrupolar interaction, that is, the coupling of the nuclear quadrupole moment with the electric field gradient of the C-2H bond. For deuterons in C-2H bonds this can lead to a splitting of about 250 kHz. As in the case of dipolar interaction, a Pake spectrum is obtained for a powder sample. The z-principal axis of the quadrupolar interaction is oriented along the bond axis which makes deuteron NMR particularly useful for studies of segmental orientations and molecular dynamics (reorientation) [1]. In sufficiently mobile, (i.e., liquid-like), systems, the anisotropy is averaged out by the isotropic thermal motions leaving only the isotropic contributions. As already stated, viscoelastic polymers represent an intermediate between the two extremes of rigid or mobile materials and the implications of this will be discussed in 14.2.3 in more detail.

14.2.2 Manipulation of Spin Interactions The rich information content of solid-state spectra makes them difficult to evaluate, in particular if more than one of the interactions discussed previously has to be taken into account. The evaluation of the spectra would often be impossible but NMR methodology provides the possibility to decouple and recouple spin interactions nearly as desired (see [11] for a comprehensive introduction). Moreover, these different information sources can be separated and correlated using the multidimensional techniques discussed subsequently. The most prominent example of a technique for decoupling or line-narrowing is magic angle spinning (MAS) (see also [1,9,11]). Here, the angular dependent part of the interactions is modulated by rapidly spinning the sample around an axis inclined at an angle Θ to the magnetic field. If the spinning axes is chosen along the so-called magicangle Θm=54.7°, the relevant scaling factor (3cos2Θ -1)/2 becomes zero and the anisotropic part of the interaction vanishes. An example using an elastomer is shown in Figure 14.1. While the static spectrum is broadened by (already motionally averaged) dipolar couplings (Figure 14.1a), these broadening effects are removed in the MAS spectrum (Figure 14.1b). Often it is necessary to manipulate the spin interactions also by pulse techniques [1,9,11]. These act on the spin operators in the corresponding Hamiltonians (see for example Equation 14.2) rather than on the geometric part. Depending on the applied pulse

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Figure 14.1 1D spectra of a typical rubber, styrene-butadiene-rubber (SBR). a) Static 1 H spectrum acquired at a Larmor frequency of 500 MHz. The dipolar coupling is motionally averaged and different lines can be distinguished although they are still broadened by the residual dipolar couplings. b) MAS spectrum of the same sample at a MAS spinning frequency of 15 kHz. The line-broadening due to anisotropic spin interactions, e.g., residual dipolar couplings, is removed Figure 14.1b reproduced with permission from R. Graf, D.E. Demco, S. Hafner and H.W. Spiess, Solid State Nuclear Magnetic Resonance, 1998, 12, 139, Figure 2b. Copyright 1998, Academic Press

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Spectroscopy of Rubbers and Rubbery Materials sequence, a given spin interaction can be switched on and off in order to discriminate between the different contributions to the desired information. For instance, 13C chemicalshift values can be determined by selectively irradiating the protons. This so-called heteronuclear dipolar decoupling removes any influence of the coupling to the protons from the 13C spectrum. Using long radio-frequency pulses on both sides with well-matched amplitudes, the flip-flop part of the heteronuclear dipolar interaction can be preserved. This can be used for the transfer of the proton polarisation to the carbon side (cross polarisation) [1,9] which is often used to artificially increase the polarisation of the dilute spins to improve the signal-to-noise ratio. Sometimes it is also applied to edit those of the dilute spins that are closely coupled to abundant spins. Undesired homonuclear spin interactions can be also suppressed using suitable multiplepulse sequences while still exploiting the information content provided by interactions that are not affected. Using a combination of MAS and pulse decoupling it is even possible to reintroduce parts of an interaction that would be averaged out by one of the manipulation techniques alone (‘recoupling’) [11]. This high flexibility of solid-state NMR enables one to fully exploit the rich information content provided by the spin interactions. It becomes particularly powerful if such experiments are combined to multidimensional NMR techniques as discussed in Section 14.3.

14.2.3 Residual Couplings and Dynamic Order Parameters In polymers and in particular in elastomers, molecular dynamics plays an important role and largely determines the mechanical properties. An investigation of segmental motions is thus of prime importance and NMR is a particularly suitable technique for this purpose. The measurement of dynamics by NMR takes place via the influence of motions on the anisotropic spin interactions. Modulation of the spin interactions by molecular motion is responsible for the relaxation of the spin system back to the equilibrium state. This socalled longitudinal (T1) relaxation has to be distinguished from transverse (or T2) relaxation that describes the dephasing of coherences under influence of residual, (i.e., partially averaged), spin interactions. Both are thus sensitive probes for molecular motions in (viscoelastic) polymers as will be briefly discussed in Section 14.2.4. If the dynamics of individual groups are of interest, the motions can be followed in great detail by twodimensional techniques on the time scale of tens of milliseconds [1]. Such investigations have proved to be extremely valuable for the investigation of polymers around the Tg. At temperatures of about 50 k above Tg, which is the more interesting regime for rubbery materials, the motions are rapid enough (tens of kHz) to effectively average the spin interactions. However, they are not fully anisotropic in high-molecular-mass polymers

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Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials and leave residual dipolar couplings as a consequence of the anisotropy [8] (see also Figure 14.1a). Restrictions of the chain motions result on a larger length scale from entanglements and crosslinks and on a shorter length scale from local chain order due to the presence of local conformations or stiffening structural elements. For characterising these effects, an autocorrelation function can be defined which is of the form S(t) = < (3cos 2 θ (t) - 1) > / 2. Here, θ(τ) is the instantaneous angle relating a vector fixed to a structural element and represents the average over segmental orientation and time. This function will not immediately decay to zero but to a certain level which can be considered as a dynamic order parameter of segmental motion (see Figure 14.2). On a longer time scale isotropisation can occur if no permanent crosslinks are present and S -> 0. Since this plateau represents the interesting regime for characterising viscoelastic properties of a polymer, we will in the following concentrate on multidimensional NMR techniques for measuring site-selectivity the order parameter or, correspondingly, residual couplings. First however, a brief overview on 1D NMR investigations serving a similar purpose will

Figure 14.2 Scheme of a two-step decay of the correlation function S of a unit undergoing anisotropic motion before finally full isotropisation is achieved. The dashed line indicates the case of a one-step decay by isotropic motion. For the study of viscoelastic polymers the intermediate plateau that reflects residual couplings is the most interesting part

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Spectroscopy of Rubbers and Rubbery Materials be given, since they form the basis for the more involved multidimensional NMR studies. A more detailed review for the particular case of crosslinked elastomers can be found in reference [12].

14.2.4 One-dimensional NMR Studies of Molecular Motions and Dynamic Order As already stated, molecular dynamics in polymers have to be probed on various time scales from relatively fast segmental motions (with correlation times of the order of 10-10 s) down to slow motions in the 10-3 s range or even slower. Dynamics in the range of around 10-10 s can be investigated by T1 measurements which are sensitive to motions around the Larmor frequency. Particularly useful for such investigations is the field cycling technique, where the magnetic field and thus also the Larmor frequency ω can be varied over several orders of magnitudes down to the 10 kHz regime. By acquiring T1(ω), molecular dynamics can be probed over the corresponding range [10] (and references therein). Slow motions can be studied using so-called longitudinal relaxation in the rotating frame (or T1ρ). In these experiments a radio frequency field (‘spin lock pulse’) with field strength of around 10-100 kHz plays the same role as the (much stronger) main magnetic field does for T1 [10]. For example, the changes in the phase structure under mechanical deformation of thermoplastic elastomers have been investigated by T1ρ measurements of 13 C nuclei under MAS [13]. A technique for the study of ultra slow molecular motions in viscoelastic materials is based on the measurement of the stimulated and primary echoes of a three-pulse experiment [10]. The ratio of the echo amplitudes depends on the characteristic correlation times, the second moment of the dipolar fluctuations and the average dipolar constant which is proportional to the order parameter of the system [14,15]. This so-called ‘dipolarcorrelation effect’ (DCE) was used for the study of ultra slow molecular dynamics and segmental order in polymer melts and networks [16,17]. The dipolar correlations in SBR networks were found to exist on a time scale exceeding 300 ms [16] and the mean squared fluctuations of the dipolar coupling constant depends on the number of segments between crosslinks N, according to N-0.75. This was interpreted by assuming that the meshes of the network are more stretched than expected for Gaussian random coils [17]. Transverse relaxation (T2) and line shape analysis has also been the topic of many studies of molecular dynamics in viscoelastic media [12]. It was found to be indicative of

526

Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials heterogeneities of molecular motions [8,18]. Transverse relaxation curves therefore often turn out to be composed of a superposition of multiexponential and Gaussian decays [8,19]. Measurement of residual dipolar couplings, molecular motions and other parameters have been performed by this relaxation technique on different nuclei like 1H, 13 C, 15N, 31P and 29Si and could be related to material properties such as the crosslink density [8,19]. In order to separate the liquid-like and solid-like contributions of the T2 decay, a linear combination of Hahn echoes and solid echoes have been used [20,21]. While Hahn spin-echo amplitudes are encoded by irreversible dephasing due to molecular motions and by reversible residual dipolar couplings, only the first effect encodes the amplitudes of solid echo (actually mixed echo) to first approximation [8,18]. The appropriate combination of these spin echoes thus gives access to residual dipolar couplings and the fluctuation rate [20,21]. The method was applied to measure the degree of segmental order induced in natural rubber upon stretching [20]. A powerful technique for the study of orientation and dynamics in viscoelastic media is line shape analysis in deuteron NMR spectroscopy [1]. For example, the average orientation of chain segments in elastomer networks upon macroscopic strain can be determined by this technique [22-31]. For a non-deformed rubber, a single resonance line in the deuterium NMR spectrum is observed [26] while the spectrum splits into a well-defined doublet structure under uniaxial deformation. It was shown that the usual network constraint on the end-to-end vector determines the deuterium line shape under deformation, while the interchain (excluded volume) interactions lead to splitting [2631]. Deuterium NMR is thus able to monitor the average segmental orientation due to the crosslinks and mean field separately [31]. The network structure of unfilled and filled elastomers was probed by the quadrupolar splitting in 2H solid-echo spectra of uniaxially strained samples [23,32] (see Figure 14.3). The local chain order at a given elongation is larger by a factor of 1.5-2 in the filled system. A decrease of local chain mobility in the absorption layer is observed under stress. The same method was applied to investigate molecular dynamic in thermoplastic elastomer based on hydrogen bonding complexes [33,34]. The 1D NMR techniques described above already allow a detailed investigation of polymer dynamics but they are mostly not selective in the sense that they do not provide information on the averaging of particular couplings. For the interpretation of some of the above results therefore a representative spin-pair along the chain was assumed, thus neglecting local site-specific motions as well as the geometry of the bonds. This is where multidimensional NMR techniques come into play.

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Figure 14.3 Deuterium solid-echo spectra of unfilled (a) and filled (b) poly[dimethylsiloxane] networks at 305 K with and without mechanical stress as given by the parameter λ [23]. This example demonstrates the sensitivity of the NMR lineshape and thus of the spin interactions to internal and external conditions Reproduced with permission from V.M. Litvinov and H.W. Spiess, Die Makromolekulare Chemie, 1992, 193, 1181, Figure 8. Copyright 1992, Wiley-VCH

Figure 14.4 (a) Principle of 2D NMR spectroscopy. The coherences are excited by a preparation sequence and evolve for a time t1. In the following mixing time the system is allowed to undergo changes after which the resulting signal is read out (direct detection). By incrementing the evolution time t1 in successive experiments, a 2D data set is produced. Fourier transformation provides the 2D NMR spectrum. (b) Simple example for the general scheme in form of a three-pulse sequence. The delays t1, t2 and tm correspond to the evolution, detection and mixing periods, respectively 528


14.3 Multidimensional NMR Spectroscopy of Viscoelastic Materials

14.3.1 Principle of Multidimensional NMR In 2D NMR spectroscopy, a two-dimensional data set is acquired as a function of two time variables t1 and t2 as shown schematically in Figure 14.4 [1, 7]. Figure 14.4a shows the general case while the three pulse sequence of Figure 14.4b represents a typical example. A 2D NMR experiment usually is divided in several time periods that follow each other. The first is the so-called preparation period in which coherences are excited by a suitable pulse sequence, which in the simplest case is only one radio-frequency pulse (see Figure 14.4b). Unlike 1D NMR spectroscopy, the excited signal is not directly acquired but is allowed to evolve in the so-called evolution period under influence of the relevant spin interactions. The evolution time t1 is incremented in subsequent experiments and provides the first time dimension of the experiment. After the evolution period (and an optional mixing time), the remaining signal is directly detected in the detection period for each time increment thus generating a 2D data set. Two-dimensional Fourier transformation then gives the 2D spectrum. Optionally a so-called mixing period of length tm can be inserted between the evolution and detection periods. During tm changes in the system can occur, for instance, by molecular motions, the action of spin interactions or by spin manipulation (for instance cross-polarisation to another nucleus). The different aspects of 2D NMR spectroscopy are reflected in the different variants that can be distinguished. One variant, separation spectroscopy, is used to separate different interactions taking advantage of the spin manipulation techniques [1,7]. For instance, during the evolution period the spin manipulation can be made such that only the isotropic chemical shift is acquired while in the detection period the full spectrum is acquired. This leaves the (motionally averaged) anisotropy to be studied site-selectively. Other 2D NMR techniques, so-called correlation techniques [1,7], aim at obtaining new information by correlating different interactions. For example, the 13C chemical shift anisotropy can be correlated with the heteronuclear dipolar powder pattern in order to obtain information on the relative orientation of the two tensors [7]. Considering the manifold of spin manipulation techniques there is a wealth of such 2D NMR techniques that can be derived for different purposes. Finally, introducing a mixing time tm, 2D exchange spectroscopy can be performed. The most important application of such exchange techniques with respect to polymer investigations is the study of slow molecular dynamics [1]. In these experiments, reorientations due to molecular dynamics are allowed to take place during the mixing time tm and lead to characteristic off-diagonal patterns in the resulting 2D spectra. If the mixing time is increased in a series of 2D experiments, slow dynamics in the range of milliseconds to seconds can be investigated in detail [1].

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14.3.2 Two-dimensional NMR Techniques and Applications to Viscoelastic Polymers The investigation of molecular dynamics and local segmental orientation by 2D NMR spectroscopy has a number of advantages [1,7] particularly when dealing with complicated spin systems such as viscoelastic polymers. Next, examples of such investigations will be discussed, concentrating on techniques that are particularly sensitive to the viscoelastic regime (see Figure 14.2), that is, NMR techniques for measuring residual dipolar couplings and dynamic order parameters. Readers interested in a more general overview of multidimensional NMR of polymers are referred to [1].

14.3.2.1 Dynamic Order Parameters from 2D 1H Magnetisation Exchange Spectroscopy Two-dimensional 1H magnetisation-exchange spectroscopy and its reduced one-dimensional variant have been used to probe residual dipole-dipole couplings between different functional groups of an SBR sample, namely between the CH- and CH2-groups of the butadiene units [35,36]. The intergroup residual dipolar couplings can then be correlated with the shear modulus which is an independent measure of the crosslink density [36]. Hence, the dynamic local-order parameter associated with the partially averaged dipolar coupling can be evaluated taking the average over the relevant dynamic processes.

• Principle of the method The basic radio frequency pulse sequence is depicted in Figure 14.4.b. Following the general scheme of 2D NMR spectroscopy, an evolution period of duration t1 follows the excitation pulse. This interval ends with a flip-back pulse and the system is allowed to exchange during the mixing period tm. After a third pulse the remaining signal is acquired during the detection interval t2 [36]. The intensity of the crosspeaks in the resulting 2D spectra then reflects the degree of exchange between the corresponding groups. Since full 2D NMR spectroscopy is time consuming, for quantitative evaluations a reduced onedimensional form has also been used. In this case the time t1 is fixed and adjusted such that it acts as a chemical-shift filter [1] for one of the two butadiene lines in the static SBR spectrum. The reappearance of the filtered line with increasing mixing time is then a measure of the magnetisation exchange between the protons of the CH and CH2 groups. Note, that unlike the 500 MHz spectrum in Figure 14.1, the styrene line and the CH line could not be distinguished in these experiments so that only the CH and CH2 lines of the butadiene units are resolved. They are separated by about 3.35 ppm, corresponding to 1 kHz at the Larmor frequency of 300 MHz.

530

Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials The theory needed for the evaluation of the exchange process was also developed [36]. In the limit of short mixing times tm, the initial decay rates for the CH line were found to depend only on the intergroup residual dipolar coupling D CH − CH 2 according to I z (t m )

= 1−

1 2

(D

CH − CH 2

)

2

t 2m

(14.3)

where I z ( t m ) is the observable NMR for the CH protons (for the CH2 protons the factor 1/2 must be substituted by 1/4 in Equation 14.3). The symbol (...) represents the statistical ensemble average of the space part of the dipolar coupling in a disordered elastomer. Note, that from such measurements, the intergroup dipolar coupling D CH − CH 2 can be determined which is largely parallel to the segment axis. That is, it is much more informative for the detection of chain order and chain motion than intragroup couplings would be. For short mixing times, the measurable exchange of longitudinal proton magnetisation is predominantly determined by the protons belonging to nearest-neighbour CH- and CH2-groups. At longer mixing times on the other hand, a much larger spin system along the polymer chain must be considered, in principle. In reference [37], however, it was shown that even in this case the intergroup residual dipolar coupling can be measured from the initial-time regime of the magnetisation-exchange decay curve.

• Two-dimensional 1H magnetisation-exchange processes in crosslinked SBR From the point of view of molecular motion, crosslinked elastomers are highly heterogeneous systems and a solid-like behaviour is present together with slow and fast motions of both, chain segments and functional groups. Hence, the magnetisationexchange process is also expected to be heterogeneous, that is, driven by different mechanisms. To investigate this heterogeneity, a series of static 2D NMR experiments were performed on SBR by increasing mixing time [36]. As already discussed previously, the aromatic and the olefinic protons are not resolved under the applied experimental conditions as can be seen in the 2D spectra shown in Figure 14.5. A remarkable feature of the diagonal 2D spectrum for short mixing times is, that the width of the diagonal peaks perpendicular to the (ω1,ω2)-diagonal is much smaller than the widths along the diagonal. This indicates, that the 1 H NMR spectrum is heterogeneously broadened [36]. The width of each line along the diagonal reflects both,

531


Figure 14.5 Two-dimensional 1H magnetisation-exchange spectra of a SBR sample recorded with the three-pulse sequence of Figure 14.4b. Different mixing times have been used: tm = 2.5 ms, (a) 5 ms, (b) 15 ms, (c) and 30 ms (d). The 2D surface representation for tm=30 ms in (e) shows that all cross-peaks are positive [36] Reproduced with permission from D.E. Demco, S. Hafner, C. Fülber, R. Graf and H.W. Spiess, Journal of Chemical Physics, 1996, 105, 11285, Figure 4. Copyright 1996, American Institute of Physics

532

Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials the solid- and the liquid-like contributions, whereas the width perpendicular to the diagonal reflects selectively the liquid-like linewidth only. This feature is very pronounced for the CH diagonal peak and less for the CH2 diagonal peak because of the strong proton-proton coupling within the methylene group. A similar result was already found in a 13C magnetisation-exchange experiment on natural rubber aimed to detect chain diffusion [38]. While no chain diffusion was detected, the lines were also found to be spread out along the diagonal. In the 1H magnetisation-exchange experiment on SBR not much cross-peak signal is found for mixing times shorter than 2.5 ms (see Figure 14.5a). This indicates, that within this time scale, the CH and CH2 functional groups can be considered in reasonable approximation to be isolated. The inter-group residual dipolar coupling leads to magnetisation exchange for longer mixing times. Intense positive cross-peaks are thus present for tm = 30 ms (see Figures 14.5d and 14.5e). Such positive cross-peaks can arise for two reasons: from the exchange process mediated by residual dipolar couplings and from cross-relaxation induced by segmental motions with correlation times τc longer than ω0-1 [7] (see also Section 14.3.2.2).

• Proton intergroup residual dipolar couplings The residual dipolar couplings between the protons of the CH group and the CH2 group were determined from the tm-dependence of the peak intensities. The decay of the longitudinal magnetisation M(tm) of the diagonal signals (normalised to the value M0 that corresponds to zero mixing time) is recorded for short mixing times for the CH group. According to Equation 14.3, the magnetisation-exchange dynamics should show an initial quadratic dependence on tm. This dependence is indeed found in the experiments and is shown in Figure 14.6a for SBR as an example. From the slope of the initial magnetisation decay, the values of the effective residual CH − CH 2 dipolar couplings Deff ≡

(

CH − CH 2 Deff

)

2

1/2

could be evaluated (see Figure 14.6a) and

it could be shown that these couplings scale with the crosslink density [36]. The correlation CH − CH 2 of the coupling constant Deff with the shear modulus G as an independent measure of the crosslink density demonstrates this behaviour (Figure 14.6b). Note, however, that the line in Figure 14.6b does not cross the origin but provides a significant value for CH − CH 2 Deff at G=0. This is not in contradiction with the applied model, but reflects the influence of local chain order (physical crosslinks) [36]. From extrapolating to G=0 in Figure 14.6b, the inter-group residual coupling of an uncrosslinked SBR melt was CH − CH 2 determined to be Deff ≈533 Hz. Using average internuclear distances and bond angles

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Figure 14.6 (a) The initial decays of the proton magnetisation for the exchange process CH® CH2. The initial part of the magnetisation-exchange data (starting with tm=100 s) shows the predicted dependence in t 2m . The effective CH-CH2 intergroup dipolar CH − CH 2 coupling constant Deff can be evaluated by fitting the theoretical curve. (b) CH − CH 2 Deff versus the shear modulus G for a series of differently crosslinked SBR samples [36] Reproduced with permission from D.E. Demco, S. Hafner, C. Fülber, R. Graf and H.W. Spiess, Journal of Chemical Physics, 1996, 105, 11285, Figure 6. Copyright 1996, American Institute of Physics

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Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials for SBR, one can extract the value for the dynamic-order parameter P2

H − H2

r CH − CH 2 − CH 2 = DCH P2 CH-CH2 intergroup linkage rCH − CH 2 according to Deff rigid

H − H2

of the . This

≈ 0.1 and suggests an even higher value for the dynamic-order parameter r r for the corresponding carbon-carbon bond P2 C − C ≈ 0.13, since rCH − CH 2 and rC − C form

yields P2

H − H2

an angle of 24°. Thus, even in polymer melts the chain motion can be largely anisotropic, as shown previously for poly(methacrylates) [39]. Magnetisation-exchange NMR spectroscopy offers a convenient means for measuring the residual dipolar couplings between functional groups along the polymer chains. The intergroup dipolar coupling is remarkably high even in uncrosslinked melts representing relatively high local segmental order. In crosslinked samples, it scales with the crosslink density as predicted by the scale-invariant model of residual dipolar couplings [8], emphasising its relation to viscoelastic properties of the elastomers.

14.3.2.2 Molecular Motions from 2D NOESY-MAS Spectroscopy Previously the coherent magnetisation-exchange process by 1H non-vanishing average dipolar couplings was discussed. Under fast MAS conditions, however, the residual couplings responsible for the exchange are largely refocused for full rotor periods and cross-relaxation by the fluctuating part of the dipolar interaction predominates. This socalled Nuclear Overhauser effect (NOE) [7] takes place if the spin system is not in internal equilibrium. It corresponds to relaxation between dipolar coupled nuclei such as 13C and 1 H or between different sites (lines) of the same type of nucleus. NOE spectroscopy (NOESY) is well-known in liquid-state NMR where it is one of the standard 2D techniques to elucidate and assign structures of macromolecules in solution [7]. Following the success of the technique in the solution-state [40], NOESY has been also applied to polymers or viscoelastic materials [35,36,41-47]. As in solution experiments, one takes advantage of the fluctuating part of the dipolar coupling to extract useful structural and dynamical information. Using this method, NOE factors have been measured in polyisoprene over a range of temperatures with static 13C NOE spectroscopy [41]. A molecular-weight independent change of regime was observed at around 60 °C for the backbone motion reflecting a loss of motional cooperativity with increasing temperature. Also the temperature-independent correlation time of the internal rotation of the methyl group could be characterised. 13

C NOE spectroscopy under MAS was used for probing polymer miscibility in polymer blends, polystyrene/polyvinyl methyl ether (PS/PVME) [42]. This study takes advantage of the fact that crosspeaks appear only between spins that are neighbours of each other,

535

Spectroscopy of Rubbers and Rubbery Materials thus establishing NOE as a probe for the degree of mixture on the molecular level. Additional information on the molecular structure of the blend could be obtained from the NOE growth rates. The results suggest that there exists a specific interaction between the phenyl ring of the PS and the PVME methyl group. In reference [43] the same technique was applied for investigating methyl groups as a source of cross-relaxation in solid polymers such as polycarbonate or polystyrene. Static 1H 2D NOE spectroscopy was applied in a first experiment showing that the technique can be used to measure inter-chain interactions [44]. This work was then continued by applying the technique under MAS to investigate the inter-molecular interactions responsible for the miscibility in polybutadiene/polyisoprene blends above the Tg [45]. It was shown that intermolecular association can be probed by this technique and the results reveal the existence of weak intermolecular interactions between the polyisoprene methyl group and the vinyl side chain of the polybutadiene.

• 1H NOESY MAS applied to the investigation of molecular motions Recently 1H NOESY MAS was used to study molecular motions in technically relevant materials such as rubbers [46, 47]. For the evaluation of these parameters, it is necessary to understand the cross-relaxation process in the presence of anisotropic motions and under sample spinning. Such a treatment is provided in [47] and the cross-relaxation rates were found to weakly depend on fast motions in the Larmor-frequency range and strongly on slow motions of the order of the spinning frequency νR. Explicit expressions for the νR dependent cross-relaxation rates were derived for different motional models. Examples explicitly discussed were based on a heterogeneous distribution of correlation times [1,8,48] or on a multi-step process in the most simple case assuming a bimodal distribution of correlation times [49-51]. The derived relationships were tested experimentally on a crosslinking series of SBR and the cross-relaxation was studied as a function of the rotor frequency νR. As an illustration of such measurements, Figure 14.7a shows a surface representation of the gradient-selected 2D NOESY spectrum of an SBR sample acquired with a mixing time of 2.7 ms at a rotor-frequency of 8 kHz. The lines corresponding to the aliphatic, the olefinic and the aromatic protons are well resolved and can be assigned as indicated. Already at short mixing times pronounced cross-peaks are visible, in particular between the olefinic CH and the CH2 group. Figure 14.7b shows the decay curve of the diagonal peaks for mixing times up to 3 ms. Unlike the magnetisation-exchange case [36,37], the decay is found to be approximately linear in tm in the short mixing time regime and exponential or biexponential for long

536


Figure 14.7 (a) 2D 1H-NOESY surface spectrum of SBR acquired for a mixing time of 2.7 ms and under 8 kHz MAS [47]. Already at this short mixing time pronounced crosspeaks are visible. (b) Short time decay of the crosspeak intensity evaluated from a series of 2D spectra. Such curves can be analysed, providing information on internuclear distances and molecular dynamics (correlation times) Reproduced with permission from T. Fritzhanns, D.E. Demco, S. Hafner and H.W. Spiess, Molecular Physics, 1999, 97, 931, Figures 2c and 3a. Copyright 1999, Taylor & Francis

537

Spectroscopy of Rubbers and Rubbery Materials times [47]. A similar curve can be plotted for the increase of the cross-peaks. To investigate the effect of the sample spinning on the cross-relaxation rates, series of 2D NOESY MAS spectra were acquired as a function of the rotor frequency. The rates evaluated from these series were found to decrease linearly with νR which is an indication that crossrelaxation in elastomers is dominated by slow motions in the 10 kHz regime rather than by motions in the Larmor-frequency range as in case of liquids. The cross-relaxation rates at room temperature were found to depend only moderately on the crosslink density which is another indication that cross-relaxation is dominated by the α process. The experimental rates were analysed in terms of the explicit expressions derived by the theoretical treatment [47]. As expected for a statistical copolymer, segmental motions in the SBR samples cannot be described by a single correlation time [1]. From T1 data that were also measured the correlation time should be in the range of 10-8 s while the νR dependence of the cross-relaxation rates requires a correlation time of 10-5 s. So it is clear that at least two different correlation times are necessary to account for the experimental findings. The relaxation dispersion data T1(ω) (ω being the Larmor frequency) [52], however, show no discontinuity that would indicate the presence of two distinct motional processes. Thus, the data were analysed in terms of a broad distribution of correlation times for almost isotropic segmental motions (α relaxation). With a simple log-Gaussian distribution function of reasonable parameters [47,48] it was found that one can account for both, the T1 values and the νR-dependent cross-relaxation rates. From the centre of gravity of the distribution, the Tg could be estimated as Tg = –43 °C using the well-known Williams-Landel-Ferry (WLF) equation [1]. This value compares favourably with the known value of Tg = –53 °C for uncrosslinked SBR-1500.

14.3.2.3 Selective Residual Dipolar Couplings by 1H Multiple-Quantum NMR Spectroscopy Multiple-quantum (MQ) NMR spectroscopy is well-established for structural studies of liquids and highly mobile solutes in liquid crystals [7,53,54]. In recent years there has been a sustained effort to obtain homonuclear [55-63] and heteronuclear [64-66] high-resolution MQ NMR spectra also for organic solids using fast MAS to increase resolution and sensitivity. Such MQ spectra proved to be valuable tools for determining dipolar connectivities between spin-1/2 nuclei [56-58, 60-63]. More quantitatively, dipolar couplings, internuclear distances and molecular torsion angles can be measured by these techniques [56-63]. For viscoelastic materials it was recently shown, that 1H high-resolution MQ NMR spectroscopy offers the possibility to measure site-selective residual dipolar couplings between all resolved protons [60, 61]. Thus the MQ technique is an attractive tool for studying structure and dynamics in polymer melts [60] and elastomers [61-63].

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Figure 14.8 (a) General scheme of two-dimensional multiple-quantum (MQ) spectroscopy. During both intervals of free precession (MQ coherences during t1 and SQ coherences during t2) dipolar decoupling can be achieved by MAS. Possible pulse sequences used for excitation/reconversion of MQ coherences are: (b) broadband BABA [57] and (c) C7 [67] Reproduced with permission from R. Graf, D.E. Demco, S. Hafner and H.W. Spiess, Solid State Nuclear Magnetic Resonance, 1998, 12, 139, Figure 1. Copyright 1998, Academic Press

• Principles of the method The MQ experiment [53, 54] follows the general scheme of 2D NMR spectroscopy is shown in Figure 14.8a. In the beginning, MQ coherences are excited (preparation period) and evolve for an evolution time t1. At the end of the evolution period, the MQ coherences are reconverted to detectable single-quantum (SQ) coherences (mixing period) which then are detected.

539

Spectroscopy of Rubbers and Rubbery Materials The experiment is performed under fast MAS conditions, that is, the rotor frequency ωR is larger than the anisotropic spin interactions. Hence, an efficient averaging of the anisotropic spin interactions takes place during the evolution and detection periods leading to well-resolved, nearly liquid-like NMR spectra. On the other hand, during the excitation and reconversion periods of the experiment dipolar interactions are required for the generation (and reconversion) of the MQ coherences. The relevant parts of the dipolar Hamiltonian are thus reintroduced by broadband dipolar-recoupling pulse sequences such as back-to-back (BABA) [57] (Figure 14.8b), or C7 [67] (Figure 14.8c). Note, that only those spins contribute to the double quantum (DQ) signal that are relatively strongly coupled to each other.

• Dipolar connectivities from the high-resolution 1H DQ MAS spectrum Figure 14.9b shows the 1H DQ spectrum of a crosslinked SBR sample which has been acquired at a spinning frequency of 10 kHz [61].

Figure 14.9 Figure 14.9 1H DQ MAS spectrum of a SBR sample with a low value of crosslink density (0.8 phr-0.8 phr sulfur-accelerator) as shown in b [61]. The rotor frequency was 10 kHz and the t1 increment was 15 μs. The diagonal peaks and the cross peaks of the functional groups are assigned as indicated (compare (a) and (b)) Reproduced with permission from R. Graf, D.E. Demco, S. Hafner and H.W. Spiess, Solid State Nuclear Magnetic Resonance, 1998, 12, 139, Figure 3. Copyright 1998, Academic Press

540

Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials On top of the 2D spectrum, the single-quantum projection is shown, which corresponds to the MAS spectrum in Figure 14.1 apart from some deviations caused by DQ-filter effects. The peaks in the 2D DQ spectrum correspond to double-quantum coherences between two spins which must be relatively close neighbours in space in order to contribute significantly to the peak intensity as follows from the strong distance sensitivity of the dipolar coupling, Equation 14.2. From the existence of the corresponding peaks therefore through-space dipolar connectivities can be easily established. The assignment of the peaks is shown along the ω1 dimension by pairs of letters a-f (see Figure 14.9a) which indicate the functional groups that are involved in the generation of the corresponding DQ peaks. From a simple qualitative inspection of the DQ spectrum it can be seen that there are dipolar connectivities between practically all functional groups. The strongest DQ signals are found between protons of the polybutadiene groups, but also DQ peaks of considerable intensity are visible between the polybutadiene protons and the aromatic protons of the styrene units. This indicates a good mixture of the different functional groups on the nanometer length scale as is found for instance in the case of a statistical copolymer.

• Site-selective double-quantum build-up curves and correlation with crosslink density A more quantitative evaluation of the DQ peak intensities is possible by acquiring the experimental DQ build-up curves for each of the peaks performing a series of DQ experiments with increasing excitation time. As an example, the build-up curves for the three main contributions are shown in Figure 14.10 for two samples with different crosslink densities [61]. A polynomial fit of the experimental DQ built-up curves using a relationship derived in the spin-pair approximation allows the ratio of the corresponding residual dipolar

(

)(

)(

)

− CH 2 2 − CH couplings to be determined, i.e., D CH 2 SCH : D CH 2 − CH SCH : D CH − CH SCH , s s s

where D represents a pre-averaged dipolar coupling constant and Sijs the corresponding 2 − CH scaled dynamic-order parameter [60, 61]. The order parameter SCH (calibrated using s CH 2 − CH a 1D MAS sideband pattern) is found to be close to the value Ss = 0.1 that has ij

been estimated for SBR samples by the magnetisation-exchange experiments described above [36]. The order parameter corresponding to the CH-CH coupling is even higher [61]. This coupling provides the best measure for the chain dynamics since it is predominantly aligned along the segmental axis. For investigating a series of samples with different crosslink densities, however, the procedure described previously is very time-consuming. In this case, a somewhat less

541


Figure 14.10 Site-selective 1H DQ build-up curves for low (a) and high (b) crosslinked SBR samples [61], respectively. The DQ signals correspond to the -CH2- (■), CH2-CH- (▲) and CH=CH coherences (●). The vertical dashed lines mark the excitation times for the maximum signal for the methylene protons, which differs for the two cases. More quantitatively, the build-up curves can be evaluated with respect to residual dipolar couplings Reproduced with permission from R. Graf, D.E. Demco, S. Hafner and H.W. Spiess, Solid State Nuclear Magnetic Resonance, 1998, 12, 139, Figure 4. Copyright 1998, Academic Press

542

Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials detailed 1D measurement of the integral DQ-filtered signal can be used to establish the relationship between the total DQ intensity and the crosslink density. The square root of the DQ intensity is found to be an approximately linear function of the crosslink density but, more important, the topological constraints that are already present in the melt are found to dominate the residual couplings (see also [60]). Thus the technique is not so much a sensitive measure of the crosslink density but is rather a sensitive and selective probe for local dynamic chain order. The strongest residual dipolar couplings were also edited in a non-rotating crosslinked poly(isoprene) series by exciting DQ and triple-quantum coherences in the short time regime [62,63]. From this, the dynamic order parameters of the methylene and methyl groups were estimated and correlated with the crosslink density. Essentially the same behaviour was found as for SBR.

14.3.2.4 Residual Local Dipolar Fields by Heteronuclear Correlation Experiments The NMR spectrum of rare spins, (e.g., 13C), contains extremely useful information because of the usually high chemical-shift dispersion and the possibility to easily eliminate undesired spin interactions. It would thus be advantageous to combine this high-resolution dimension with 1H spectroscopy in the form of a heteronuclear correlation experiment. Next, the so-called wide-line separation (WISE) experiment [1,68-71], in which 1H broadline information is correlated with the 13C spectroscopic information will be discussed. There are however also other possibilities for including heteronuclear information in a 2D experiment, for instance, 2D heteronuclear J-resolved spectroscopy which was applied for the investigation of filled natural rubber [38]. It was concluded that there must be a high degree of motion to allow the scalar 13C-1H couplings to be revealed by MAS alone. The 13C linewidth was found to be determined by susceptibility effects due to the presence of the filler. While the (heteronuclear and homonuclear) residual dipolar couplings have to be eliminated in such J-resolved experiments, they are the main source of information in the WISE experiment described in the next section. For the investigation of viscoelastic materials the experiment thus must be performed under static conditions or under slow MAS.

• Principles of the method The 2D WISE experiment [68, 71] is based on the simple pulse scheme presented in Figure 14.11.

543


Figure 14.11 Basic 2D pulse sequence used for the measurement of residual dipolar local fields. After the excitation pulse the spins are allowed to evolve for some time t1 (indirect dimension) under influence of the relevant spin interactions before crosspolarisation takes place. The direct detection during time t2 then takes place on the 13 C side typically under proton dipolar decoupling (DD). The basic scheme can be extended by various spin manipulation techniques (not shown) during time t1. For instance, the heteronuclear dipolar contribution can be removed by a decoupling pulse on the 13C side [73]

After the excitation of the protons by a 90° pulse, a variable period t1 follows during which a decay of the magnetisation by homonuclear and heteronuclear dipolar couplings takes place. Optionally, the heteronuclear couplings can be eliminated during this time by S-spin decoupling using a suitable pulse sequence applied on the S-spin side during the t1 period [68]. At the end of this period, a rare spin magnetisation is created by crosspolarisation. The duration of the cross-polarisation process is chosen short enough that spin diffusion between protons can be neglected. Finally, the 13C signal is recorded in the t2 dimension in the presence of proton dipolar decoupling (DD). The resulting 2D data set can then be Fourier transformed in both dimensions (resulting in a 2D WISE spectrum) or only in the 13C dimension. In the latter case the information is obtained in form of the 1 H decay for each resolved 13C line (13C edited 1H relaxation). The experiment is sometimes performed under slow MAS conditions to obtain well-resolved 13C lines. For viscoelastic materials, the spinning must be very slow in order to not drastically influence the proton lines which contain the information on the heteronuclear and homonuclear residual couplings. As an alternative, one could also generate a sideband pattern by MAS and evaluate it with respect to the residual couplings (see below). In any case, the experiment

544

Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials provides a direct visualisation of the local dipolar field for functional groups with different molecular mobilities. One should, however, be aware that it is often not only the 1H-1H dipolar coupling but also the 1H-13C heteronuclear dipolar coupling that affects the magnetisation decay of those protons that are close enough to the 13C nuclei to cross polarise them [35,68,7274]. This effect takes place because the 1H-13C internuclear distance is smaller than the 1 H-1H distance, for instance, in a methylene group. The heteronuclear dipolar coupling thus often is of the same order of magnitude or even stronger than the homonuclear dipolar coupling [68,74].

• Residual heteronuclear local dipolar couplings by WISE-MAS experiments For elastomers and rubbery-like materials well above the Tg, the high molecular mobility reduces the dipolar couplings dramatically. The WISE experiment allows one to investigate site-selectively residual dipolar interactions and thus molecular dynamics by editing the corresponding proton slices of the 2D data set. In references [35,74] 13C-edited 1H transverse relaxation was investigated in a SBR crosslinking series. A short contact time was used to avoid 1H spin diffusion. As a consequence of this, only 1H atoms directly bonded to 13C atoms are observed and the 1 H transverse relaxation is thus found to be mainly governed by 1H-13C heteronuclear couplings which makes analysis simple. The 13C-edited 1H transverse relaxation could be fitted with only one adjustable parameter, the effective number of statistical segments [35]. This was invoked as a justification for the heterogeneous model used to describe transverse relaxation. Moreover, the effect of crosslinking could be investigated and it was shown that it affects the dynamics of all functional groups to the same extent [74]. A similar experiment has been performed [68] but under MAS conditions. For a series of crosslinked natural rubber samples (A - F1), 13C edited 1H spinning sidebands have been extracted from the 2D spectrum. These sideband pattern are encoded by the residual dipolar couplings of the corresponding functional groups and are presented in Figure 14.12. For the experimental conditions given, the spectra are found to be dominated by heteronuclear dipolar interactions which leads to relatively narrow, well-separated dipolar spinning sidebands. The heteronuclear dipolar couplings could be evaluated by simulating the spectra on this basis (see right-hand column in Figure 14.12). They were found to be between 0.9 kHz and 1.5 kHz for the samples of the series (the inter-crosslink masses Mc were between 6700-11000 g/mol). A practically identical envelope of the spinning-

545


Figure 14.12 Slices from a 2D experiment corresponding to the pulse sequence of Figure 14.11 performed on a series of crosslinked natural rubber samples A- F1 (defined in [68]) under MAS. The slices reflecting proton sideband pattern for the different functional groups are encoded by the 1H residual dipolar couplings. The distinct features of dipolar slices prove that the different functional groups may be considered as relatively isolated groups of spins on the time scale of the evolution and cross polarisation Reproduced with permission from C. Malveau, P. Tekely and D. Canet, Solid State NMR, 1997, 7, 271, Figure 3. Copyright 1997, Academic Press

546

Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials sidebands intensities in the indirect dimension is found for all functional groups. This indicates that the effective 1H- 13C vectors of different segments are affected in approximately the same way by internal motions. Spin diffusion as an alternative explanation of this effect could be excluded by a similar experiment with additional 13C decoupling during the evolution period t1 which eliminates the encoding by the heteronuclear dipolar couplings [68]. Distinct features are then visible for the spinning sideband patterns of the CH3, CH2 and CH groups (not shown) proving that different functional groups are relatively isolated from each other on the time scale of the experiment as a consequence of the reduction of homonuclear dipolar interactions by high segmental mobility and MAS. Therefore, the indirect detection of the spinning sideband patterns via cross polarisation 13 C spectroscopy allows for selective measurements of residual dipolar interactions on different functional groups [68].

• Investigation of phase-separated systems using WISE Apart from the investigation of residual couplings described previously, WISE has also been applied for the investigation of phase-separated systems, for instance co-polymers, consisting of a rubbery phase and a more rigid phase [75-77]. While the different phases can be discriminated in the 13C dimension by their chemical composition, the dynamics of each of them can be studied by the corresponding 1H slice. A model polyurethane consisting of monodisperse hard segments (45% weight) and soft segments has been investigated [75] and it could be shown that both phases are completely separated from each other. The dynamics and morphology of a poly(isobutylene-co-pmethylstyrene) copolymer were investigated by various solid-state NMR techniques including WISE [76]. WISE was used to investigate dynamics by determining the transverse relaxation for polyisobutylene and methylstyrene blocks separately and differences between the two were found. Domain sizes were investigated by another 2D heteronuclear correlation experiment, where pulsed spin decoupling was used on the proton side and an additional time τ was introduced in order to allow spin diffusion to take place. The spin diffusion could be followed by the increase of the crosspeak intensities in the resulting 2D heteronuclear correlation spectra with increasing time τ. A domain size of the pmethylstyrene blocks of around 3 nm could be estimated. Two commercial blockcopolymers have been investigated [77] using among other techniques WISE. The WISE spectra showed that the mobilities of the polystyrene and rubber phases of the copolymers are very different and that the interfacial region between the two must be very small. The size of this region could be determined to be 2 nm using a spin diffusion experiment.

547


14.3.2.5 Deuterium NMR Studies on Thermoplastic Elastomers For the investigation of the molecular dynamics in polymers, deuteron solid-state nuclear magnetic resonance (2D-NMR) spectroscopy has been shown to be a powerful method [1]. In the field of viscoelastic polymers, segmental dynamics of poly(urethanes) has been studied intensively by 2D-NMR [78, 79]. In addition to 1D NMR spectroscopy, 2D NMR exchange spectroscopy was used to extend the time scale of molecular dynamics up to the order of milliseconds or even seconds. In combination with line-shape simulation, this technique allows one to obtain correlation times and correlation-time distributions of the molecular mobility as well as detailed information about the geometry of the motional process [1].

• Principles of the method The scheme typically used for such 2D-exchange NMR experiments corresponds to that given in Figure 14.4b, however, for deuterons, often a solid-echo sequence is used for detection instead of the last pulse. During the mixing time tm of the experiment molecular reorientations are allowed to take place. If the molecular orientation of a C-2H bond has changed due to slow molecular motions, the signal continues to evolve with a new frequency. For reorientation about a well-defined angle θ, the 2D exchange spectrum exhibits characteristic ridges in the form of ellipses that can be analysed for motion (for details about the analysis see [1]).

• Local motions and segmental orientation in supramolecular hydrogen bond assemblies In the field of viscoelastic materials, the technique has been mainly applied to thermoplastic elastomers. One interesting elastomer belonging to this class consist of polybutadiene chains functionalised by 4-(3,5-dioxo-1,2,4-triazolidin-4-yl) benzoic acid (U4A) units which act as effective junctions zones. The molecular dynamics of the phenyl rings of the U4A units has been probed by 1D and 2D 2H-NMR [80]. In this system, there are three spatially separated environments which are reflected in the mobility of the polar units. Phenyl rings which are incorporated in the structure are either rigid or undergo 180° phenyl flips. The small fraction of free functional groups move isotropically and their mobility is coupled to the dynamics of the polymer matrix. The 1D 2H-NMR spectra can be described quantitatively assuming a distribution of correlation times over 2-3 decades and the geometry of the motional processes is defined by the environment in the clusters up to the order-disorder transition temperature. 2D spectra of the model compounds (see Figure 14.13) show an elliptical exchange patterns, indicating well-defined slow 180° phenyl flips on a time scale of 100 ms up to 3 s.

548


Figure 14.13 2D deuterium exchange spectra of a thermoplastic elastomer (tBu-dU4A) [80]. Experimental results at 27 °C with mixing times of (a) tm=1 ms, (b) tm=500 ms and (c) tm=3 s indicating the presence of slow 180° phenyl flips Reproduced with permission from A. Dardin, C. Boeffel, H.W.Spiess, R. Stadler and E.T. Samulski, Acta Polymerica, 1995, 46, 291, Figure 12. Copyright 1995, Wiley-VCH

549


14.5 Conclusion Multidimensional NMR spectroscopy proves to be a powerful method to reveal structural and dynamical information at the molecular level in elastomers. Residual dipolar couplings can be measured site-selectively and correlated with the crosslink density and mechanical stress. The local segmental order and information on local molecular motions can be also obtained with newly developed 2D NMR methods. The information at the molecular level can be correlated with macroscopic properties of elastomers and provides the basis for a better design of material properties for specific applications.

Acknowledgements Useful discussions with Dr. G. Heinrich, Dr. H. Dumler (both from Continental AG, Hannover), with Professor B. Blümich, D. Canet, P. Tekely, and with our colleagues T. Fritzhanns, C. Fülber, R. Graf are gratefully acknowledged. DED thanks the Deutsche Forschungsgemeinschaft for a Mercator Guest Professor grant and SH for financial support of part of the work reviewed above.

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B.F. Chmelka, K. Schmidt-Rohr and H.W. Spiess in Nuclear Magnetic Resonance Probes of Molecular Dynamics, Ed., R. Tycko, Kluwer Academic Publishers, Dordrecht, The Netherlands, 1994, p.113.

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Multidimensional NMR Techniques for the Characterisation of Viscoelastic Materials 8.

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10. R. Kimmich, NMR: Tomography, Diffusiometry, Relaxometry, Springer, Berlin, 1997. 11. S. Hafner and H.W. Spiess, Concepts in Magnetic Resonance, 1998, 10, 99. 12. A.K. Whittaker, Annual Reports on NMR Spectroscopy, 1997, 34, 106. 13. A. Schmidt, W. S. Veeman, V. M. Litvinov and W. Gabrielse, Macromolecules, 1998, 31, 1652. 14. R. Kimmich, E. Fischer, P. Callaghan and N. Fatkullin, Journal of Magnetic Resonance A, 1995, 117, 53. 15. F. Grinberg and R. Kimmich, Journal of Chemical Physics, 1996, 105, 3301. 16. E. Fischer, F. Grinberg, R. Kimmich and S. Hafner, Journal of Chemical Physics, 1998, 109, 846. 17. F. Grinberg, M. Garbaczyk and W. Kuhn, Journal of Chemical Physics, 1999, 111, 11222. 18. J. Collignon, H. Sillescu and H.W. Spiess, Colloid and Polymer Science, 1981, 259, 220. 19. V.D. Fedotov and H. Schneider, Structure and Dynamics of Bulk Polymers by NMR Methods, Springer, Volume 21, Berlin, 1989. 20. P.T. Callaghan and E.T. Samulski, Macromolecules, 1997, 30, 113. 21. R.C. Ball, P.T. Callaghan and E.T. Samulski, Journal of Chemical Physics, 1997, 106, 7352. 22. B. Deloche, A. Dubault and D. Durand, Journal of Polymer Science: Part B Polymer Physics, 1992, 30, 1419. 23. V.M. Litvinov and H.W. Spiess, Die Makromolekulare Chemie, 1992, 193, 1181. 24. P. Sotta and B. Deloche, Journal of Chemical Physics, 1994, 100, 4591.

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Spectroscopy of Rubbers and Rubbery Materials 25. M. Klinkenberg, P. Blümler and B. Blümich, Macromolecules, 1997, 30, 1038. 26. P. Sotta, Macromolecules, 1998, 31, 3872. 27. M. Warner, P.T. Callaghan and E.T. Samulski, Macromolecules, 1997, 30, 4733. 28. P. Sotta, B. Deloche, J. Herz, A. Lapp, D. Durand and J.C. Rabadeux, Macromolecules, 1987, 20, 2769. 29. M.G. Brereton and M.E. Ries, Macromolecules, 1996, 29, 2644. 30. M.G. Brereton, Macromolecules, 1991, 24, 6160. 31. M.E. Ries, M.G. Brereton, P.K. Klein, I.M. Ward, P. Ekanayake, H. Menge and H. Schneider, Macromolecules, 1999, 32, 4961. 32. V.M. Litvinov and H.W. Spiess, Die Makromolekulare Chemie, 1991, 192, 3005. 33. A. Dardin, R. Stadler, C. Boeffel and H.W. Spiess, Die Makromolekulare Chemie, 1993, 194, 3467. 34. A. Dardin, H.W. Spiess, R. Stadler and E.T. Samulski, Polymer Gels and Networks, 1997, 5, 37. 35. P. Sotta, C. Fülber, D.E. Demco, B. Blümich and H.W. Spiess, Macromolecules, 1996, 29, 6222. 36. D.E. Demco, S. Hafner, C. Fülber, R. Graf and H.W. Spiess, Journal of Chemical Physics, 1996, 105, 11285. 37. L. Gasper, D.E. Demco and B. Blümich, Solid State Nuclear Magnetic Resonance, 1999, 14, 105. 38. A.P.M. Kentgens, W.S. Veeman and J. van Bree, Macromolecules, 1987, 20, 1234. 39. A.S. Kulik, D. Radloff and H.W. Spiess, Macromolecules, 1994, 27, 433. 40. D. Canet, Nuclear Magnetic Resonance, Concepts and Methods, John Wiley Publishers, Chichester, UK, 1996. 41. J. Denault and J. Prud’homme, Macromolecules, 1989, 22, 1307. 42. J.L. White and P. Mirau, Macromolecules, 1993, 26, 3049.

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Spectroscopy of Rubbers and Rubbery Materials 60. R. Graf, A. Heuer and H.W. Spiess, Physical Review Letters, 1998, 80, 5738. 61. R. Graf, D.E. Demco, S. Hafner and H.W. Spiess, Solid State Nuclear Magnetic Resonance, 1998, 12, 139. 62. M. Schneider, L. Gasper, D.E. Demco and B. Blümich, Journal of Chemical Physics, 1999, 111, 402. 63. M.Schneider, D.E. Demco and B. Blümich, Journal of Magnetic Resonance, 1999, 140, 432. 64. W. Sommer, J. Gottwald, D.E. Demco and H.W. Spiess, Journal of Magnetic Resonance A, 1995, 112, 131. 65. M. Hong, J.D. Gross and R.G. Griffin, Journal of Physical Chemistry B, 1997, 101, 5869. 66. K. Saalwächter, R. Graf, D.E. Demco and H.W. Spiess, Journal of Magnetic Resonance, 1999, 139, 287. 67. Y.K. Lee, N.D. Kurur, M. Helmle, O.G. Johannessen, N.C. Nielsen and M.H. Levitt, Chemical Physics Letters, 1995, 242, 304. 68. C. Malveau, P. Tekely and D. Canet, Solid State NMR, 1997, 7, 271. 69. P. Tekely, D. Nicole, J. Brondeau and J.J. Delpuech, Journal of Physical Chemistry, 1986, 90, 5608. 70. N. Zumbulyadis, Physics Review B, 1986, 33, 6495. 71. K. Schmidt-Rohr, J. Clauss and H.W. Spiess, Macromolecules, 1992, 25, 3273. 72. P. Tekely, P. Palmas and P. Mutzenhardt, Macromolecules, 1993, 26, 7363. 73. P. Palmas, P. Tekely and D. Canet, Solid State Nuclear Magnetic Resonance, 1995, 4, 105. 74. C. Fülber, D.E. Demco, O. Weintraub and B. Blümich, Macromolecular Chemistry and Physics, 1996, 197, 581. 75. H-J. Tao, D.M. Rice, W.J. MacKnight and S.L. Hsu, Macromolecules, 1995, 28, 4036. 76. J.L. White, A.J. Dias and J.R. Ashbaugh, Macromolecules, 1980, 31, 1880.

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556

Deuterium NMR in Rubbery Materials

15

Deuterium NMR in Rubbery Materials Bertrand Deloche and Paul Sotta

15.1 Introduction Deuterium NMR (2H NMR) has been used extensively to study and characterise molecular properties in various complex fluids such as liquid crystals [1, 2, 3] and in polymer systems [4]. Rubber materials are soft, elastic solids, made of mobile, flexible polymer chains (with a glass transition temperature (Tg) typically lower than 0 °C) which are linked together to form a three-dimensional network. They are characterised by a low, frequency independent elastic modulus (of the order 105 to 106 Pa) and usually by a large maximum reversible deformation (up to a few hundred per cent). Rubber elasticity is based on the properties of crosslinked polymer chains: at large spatial scales, the presence of crosslinks ensures the reversibility of the deformation, while at short scales, mobile polymer chains behave as molecular, entropic springs. There are different ways to realise rubber materials with ‘customised’ properties. The main requirements are usually: chemical and mechanical stability over a certain temperature range; a given value of the elastic modulus; a high maximum deformation and eventually, biocompatibility. In order to characterise the molecular behaviour in such materials, various techniques have been developed in the past 25 years. The aim of these investigations has generally been to relate the macroscopic, thermomechanical behaviour, to molecular properties, in order to improve the molecular design and to engineer materials with better performances. From a physical point of view, it is easy to understand that the local molecular orientation in rubber materials is related in a quite direct way to the polymer chain extension, which in turn gives the amount of (entropic) elastic energy stored in the rubber. Therefore, it has been essential to implement experimental approaches giving access to chain ordering in strained rubbers, in order to investigate the way in which the constraints are transmitted at a molecular scale.

557

Spectroscopy of Rubbers and Rubbery Materials Amongst the various techniques sensitive to the molecular behaviour, deuterium (2H) nuclear magnetic resonance (NMR) has some properties which make it particularly specific: •

2

•

2

H NMR provides a direct, absolute measurement of local chain extension, to which the modulus is directly related, as mentioned above. In principle, the full distribution of constraints in polymer chains may be obtained by analysing 2H NMR spectra. This property has been exploited to image the spatial distribution of constraints in a (nonuniformly) stretched rubber (see the corresponding chapter in this book); H NMR is sensitive to a very low degree of local anisotropy (or constraint);

• by selective deuteration, some contrast between various molecular groups or portions of polymer chains may be introduced in to the NMR response; •

2

H NMR may be sensitive to the presence of defects or heterogeneities at various spatial scales, which are often believed to play a major role in limiting or enhancing the elastic properties of the materials.

The chapter is organised as follows: The basics of 2H NMR are summarised in Section 15.2. Its potential for the characterisation of molecular motions, are reviewed simply. Some applications are listed to illustrate these potentials. They may be classified in various ways. The most direct approach consists in working with labelled (deuterated), model systems. ‘Real’ materials (designed for industrial applications and processed at a large scale) are often multicomponent, complex systems, which may be relatively ill defined at the molecular scale. Thus, working with chemically well defined, labelled materials (which, however, often have relatively poor mechanical properties by themselves) is a way to isolate and study the various parameters which play a role in rubber properties. Studies are done both in the relaxed state and in constrained (uniaxially deformed) states. This approach is illustrated in Section 15.3. Examples of studies performed in model, single component networks are presented. However, even in this case, the sensitivity of the method is such that it may detect the presence of a few percent of molecular defects. The necessity to label (deuterate) the systems under study may appear as a practical drawback. It often prevents working with real (industrially designed) materials. A way to circumvent this problem is to use labelled molecules as 2H NMR probes whose behaviour reflects that of the rubber matrix. These molecules may be solvent molecules or even polymer chains chemically identical to the rubber matrix itself. This approach is described in Section 15.4 and is illustrated by a few examples. 558

Deuterium NMR in Rubbery Materials In composite, multicomponent systems, 2H NMR is particularly suited to investigate interfacial properties. This is illustrated in filled rubbers, semicrystalline and thermoplastic elastomers, based on copolymers (Section 15.5).

15.2 2H NMR Background The basic concepts of deuterium (2H) NMR and its application to anisotropic fluids have been described in numerous references [1, 4-7]. This technique has some specificities which make it particularly relevant to investigate fluid, partially ordered systems: • The necessity to label (deuterate) the sample may be turned into an advantage, allowing investigation of specific components of a sample; • The predominant interaction for a 2H spin system is the quadrupolar interaction, which couples the electric quadrupole moment of the 2H nucleus to its electronic surrounding. This interaction is a second-rank tensor HQ which lies approximately along the C-2H bond in organic molecules. Thus, in practice, 2H nuclei may be considered to be isolated. It shows that the 2H NMR formalism is similar to that of an isolated proton pair [8]; • The quadrupolar interaction constant νQ is large, of the order 250 kHz, which makes 2 H NMR very sensitive to the degree of molecular mobility. In a solid system, the resonance spectrum spreads over a wide frequency range (± 250 kHz), with a structure related to the static orientational distribution of C-2H bonds in the system (powder spectrum). In a fluid system, the spectrum structure is affected by molecular motions. In the limit of fast molecular motions (νQτ , which correlates the residual interaction Δ0 at t = 0 and Δt at t = τ (τ may be as large as 0.5 s here) [30]. The exponential decay which is shown in Figure 15.3

Figure 15.3 The decay of the so-called quadrupolar order as a function of the time interval t2 (quadrupolar order is created by a pulse at time t1, then relaxes during a time interval t2 before being observed via a third pulse). Data from three differently crosslinked PDMS networks are presented (Mn = 9700 g.mol-1; o, tetraallyloxyethylene-crosslinked (tetrafunctional); + crosslinked with bis(allyloxy)-3[bis(allyloxymethyl]-2, 2-propylene oxide (hexafunctional); Δ tetravinylcyclosiloxanecrosslinked]. The longitudinal relaxation is plotted for comparison. Experimentally, the quadrupolar relaxation time T1Q equals about 1.33 T1Z Reproduced with permission from P. Sotta and B. Deloche, Journal of Chemical Physics, 1994, 100, 4591, Figure 4. Copyright 1994, American Institute of Physics

565

Spectroscopy of Rubbers and Rubbery Materials actually demonstrates that network chain end-to-end vectors have no slow reorientations. This is related to the fact that network junctions have fixed average positions in space (even though they may fluctuate rapidly around these average positions).

15.3.5 Uniaxially Deformed Model Networks Figure 15.4 shows 2H NMR spectra obtained in a uniaxially deformed PDMS network, for different values of the angle Ω between the applied force F and the magnetic field B0 [31, 32]. The line shape is drastically affected compared to that observed in a relaxed sample: a permanent, well-resolved doublet structure appears, both in stretched and

Figure 15.4 2H NMR spectra in a uniaxially compressed PDMS network (Mn = 10500 g.mol-1, deformation ratio λ = 0.5), for different values of the angle Ω between the applied uniaxial force and the magnetic field B0 Reproduced with permission from P. Sotta and B. Deloche, Macromolecules, 1990, 23, 1999, Figure 2. Copyright 1999, American Chemical Society

566

Deuterium NMR in Rubbery Materials compressed samples. The doublet splitting Δ, measured as a function of the angle Ω, reproduces with great accuracy the curve (3cos2 Ω-1)/2, in agreement with Equation 15.1. This demonstrates that the quadrupolar interactions are time-averaged along the direction of the applied force F, which means that the chain segment motions associated with the doublet are uniaxial, with a symmetry axis along this direction. Specifically, for Ω = 54.5°, the doublet vanishes and the spectrum is similar to the one obtained at λ = 1. At this point, it is important to emphasise the implications of this behaviour, as regards the physical origin of elasticity in these systems. Even though the way the anisotropy appears in the stretched system may seem rather trivial, it is not. The appearance of the doublet is not simply related to the induced overall anisotropy in a stretched rubber network. In classical, independent chain models of rubbers, elasticity is an entropic process only related to conformational restrictions in network chains, which are fixed at junction points [32, 33, 34]. The macroscopic constraint is transmitted through the junctions via the anisotropy induced in the static end-to-end vector distribution P(R). Due to end-toend stretching along its end-to-end vector R, each chain is oriented along R. As a result, the average quadrupolar interaction for one chain of N segments (x, y, z are the components of the reduced vector R/N1/2a, with z axis along B0) is [23]: Δ 0 ( R ) ≈ νQ

(

κ 2z2 − x2 − y2 N

)

(15.3)

κ is a numerical factor which accounts essentially for chain flexibility. The 2H NMR spectrum might be simulated by superposing the contributions of all chains, considered independently [35]. If the macroscopic deformation is transmitted affinely to crosslink junctions, P(R) may be written in the deformed network (the strain being along z): ⎡ 3λx2 ⎤ ⎡ 3λy2 ⎤ ⎡ 3z2 ⎤ P( R ) ≈ exp ⎢− ⎥ exp ⎢− ⎥ exp ⎢− 2 ⎥ 2 ⎥⎦ 2 ⎥⎦ ⎢⎣ ⎢⎣ ⎢⎣ 2λ ⎥⎦

where λ is the deformation ratio. A spectrum calculated in this way is shown in Figure 15.5, both in relaxed and elongated states. The most striking feature is that no doublet appears here in the elongated state [23], even though of course there is a non-zero, average overall orientation <S>o in the deformed network (the so-called Kuhn and Grün orientation function) [33]: S

0

= νQ

(

κ 2 λ − λ−1 N

)

The fact that classical theories fail in reproducing the existence of a doublet structure has been also pointed out recently [36, 37].

567


Figure 15.5 Spectra simulated with a ‘phantom’ network, in which each chain is elongated, and thus oriented, along its end-to-end vector. A gaussian distribution of the end-to-end vectors is assumed. a: relaxed (isotropic) state; b: uniaxially elongated state (deformation ratio λ = 2.5). Affine displacement of junctions is assumed. No doublet appears Reproduced with permission from P. Sotta and B. Deloche, Macromolecules, 1990, 23, 1999, Figure 4. Copyright 1990, American Chemical Society

In fact, the dynamic uniaxiality reflected in the observed doublet (together with its Ωdependence) has a completely different origin. It comes from interactions between chain segments, which, even though they are weak, have a collective effect in the deformed sample so that most of the segments exhibit the same uniaxial fluctuations along the applied force F, as observed experimentally. A mean field model for this effect was proposed [23]. In this model, the residual quadrupolar interaction in one network chain contains, in addition to the term Δ0 associated to chain elongation along the end-to-end vector (Equation 15.3), a second contribution, written as a mean-field term which represents a uniform, uniaxial contribution to orientation. This uniaxial contribution is responsible for the splitting of the spectrum in two lines. This may be written as:

Δ ≈ Δ 0 ( R ) + νQ u S

568

3 cos2 Ω − 1 2

(15.4)

Deuterium NMR in Rubbery Materials In the model, the uniform contribution (and thus, the doublet splitting) is proportional to the overall average orientation <S>. The interaction parameter u characterises the strength of orientational interactions between segments (0 < u < 1). Thus, for a given deformation ratio λ, the spectrum contains one constant splitting and a distribution of additional shifts, which is clearly seen in Figure 15.4. This mean-field effect is also clearly demonstrated by the observed orientation of free (unattached) chains chemically identical to network chains, which are introduced during or after network formation and free to diffuse in it (Figure 15.6) [20, 38, 39]. According

Figure 15.6 2H NMR spectra of perdeuterated free chains (Mn = 10500 g.mol-1) diffusing in a uniaxially constrained PDMS network (Mn = 10500 g.mol-1), for different elongation ratios λ. At λ = 1, the free chains probe an overall isotropic system, and the doublet vanishes. The fraction of free chains is 9% Reproduced with permission from P. Sotta, B. Deloche, J. Herz, A. Lapp, D. Durand and J.C. Rabadeux, Macromolecules, 1987, 20, 2769, Figure 2. Copyright 1987, American Chemical Society

569

Spectroscopy of Rubbers and Rubbery Materials to the previous model, the free chains are sensitive only to the uniaxial contribution in the stretched network (second term in Equation 15.4). Therefore, under identical conditions (same network characteristics, same elongation), they do present a permanent doublet with the same splitting (and the same Ω-dependence) as network chains, but without line broadening related to end-to-end chain elongation (first term in Equation 15.4). The doublet splitting is of course zero in the relaxed state (λ = 1), because the average orientation then vanishes and it increases with increasing deformation (elongation) ratios above [18, 17, 31]. In Figure 15.7, the splitting Δ is plotted as a

Figure 15.7 Variation of the splitting Δ in an elongated PDMS model network containing a small amount (9%) of free PDMS chains, as a function of (λ2-λ-1), with λ the elongation ratio. A pair of identical network/free chains systems are used: ▲: deuterated network chains (Mn = 10500 g.mol-1); o: deuterated free chains (Mn = 10500 g.mol-1). S is the chain segment order parameter Reproduced with permission from P. Sotta, B. Deloche, J. Herz, A. Lapp, D. Durand and J.C. Rabadeux, Macromolecules, 1987, 20, 2769, Figure 7. Copyright 1987, American Chemical Society

570

Deuterium NMR in Rubbery Materials function of the strain function (λ2-λ-1) for both network and free chains. Taking into account the internal motions and structure of the PDMS monomer [9, 10], the chain segment order parameter S is equal to a few 10-3 for λ≤2 [31, 39]. In the presence of the mean field effect introduced before, the overall average orientation is enhanced with respect to the hypothetical value <S>0 in the absence of interactions: S =

1 S 1− u

(15.5)

0

The factor 1/(1-u) is greater than 1 (u), whereas 2H NMR is sensitive to the local symmetry of molecular motions (the doublet is u<S>). In particular, the fact that the doublet is the same for network and guest chains has sometimes been misinterpreted. Strictly speaking, it does not prove that interchain interactions are strong, (i.e., u close to 1) [46], but that there is a mean field effect in the deformed network. This result may be extrapolated to network dangling chains as well. They exhibit the same orientational behaviour as free chains when the network is deformed [39]. Besides, it has been claimed that the doublet structure in the network itself may be dominated by the contribution from dangling chains [47]. This contribution has been quantified in recent 2H NMR experiments in PDMS end-linked networks with a controlled amount of selectively labelled dangling chains [17]. It is clear that both network and dangling chains contribute to the observed doublet [48]. The doublet appears in various other polymer matrices, which shows that the uniaxial ordering of chain segments is perhaps a quite general phenomenon. For instance, PB networks with chains deuterated in 1,4 positions (Figure 15.8) have been studied [1821]. Their behaviour is essentially the same, even though the linewidth is much broader

571

Spectroscopy of Rubbers and Rubbery Materials at room temperature than in PDMS. This is due to a chain mobility lower and a critical entanglement molecular weight smaller than in PDMS, which leads to a much higher viscosity. The increasing line width as λ increases is attributed here to the broad distribution of chain lengths between junctions, because shorter chains are more oriented than longer ones, and thus produce a larger doublet. Spectra can be fitted with a most probable distribution of chain lengths. Other systems are described in Section 15.5.

Figure 15.8 2H NMR spectra obtained in a randomly crosslinked, deuterated PB network. Precursor chain molecular weight 115700 g.mol-1, 1.2% crosslink agent, the average molecular weight between junctions is 27400 g.mol-1 (as determined by swelling experiments) or 11600 g.mol-1 (elastic measurements). The smooth curves are fits with a most probable distribution of chain lengths with number average molecular weight Mc = 11600 and with non-Gaussian chain statistics Reproduced with permission from W. Gronski, R. Stadler and M.M. Jacobi, Macromolecules, 1984, 17, 741, Figure 5. Copyright 1984, American Chemical Society

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15.3.6 Physical Origin of the Induced Orientation As mentioned in Section 15.3.5, the doublet has been interpreted by short-range interchain interactions. Numerical simulations using a Monte Carlo method have been performed to confirm this interpretation [49-51]. In these simulations, chains are generated as self avoiding walks on a cubic lattice, placed in a box. In order to represent network chains, extremities of chains are affixed to the lattice. Interactions are introduced only via shortrange excluded volume (which prevents double occupancy of a lattice site). The deformation is transmitted exclusively via the end-to-end vector distribution. The full distribution of orientation, (i.e., the distribution for the average orientation of each segment in the system), is sampled during the simulation. Figure 15.9 shows the simulated spectrum in an elongated network. It does present a doublet structure qualitatively similar to that observed experimentally, that varies linearly with (λ2-λ-1) and has the Ω-dependence in Equation 15.5 [49]. One advantage of numerical simulations is that a given interaction may be switched off easily: interchain (excluded-volume) interactions may be suppressed,

Figure 15.9 Spectrum for an elongated network (λ = 3.41, N = 51, Ω = 0°) obtained by Monte-Carlo simulation. The abscissa scale denotes the reduced interaction Δ/νQ, i.e. the orientational order parameter S. Ordinates are in linear arbitrary units Reproduced with permission from P. Sotta, Macromolecules, 1998, 31, 8417, Figure 1. Copyright 1998, American Chemical Society

573

Spectroscopy of Rubbers and Rubbery Materials and the spectrum obtained in a deformed network in this case does not present a doublet [49]. These simulations demonstrate that a purely isotropic interaction (excluded volume) can generate a uniaxial, (i.e., collective), orientational mean-field in a uniaxially stretched system of chains. A model, based on translational entropy, has been proposed [50]. Other interpretations have been given. It was suggested that junction fluctuations may give rise to such a uniaxial contribution in stretched samples [52]. However, this does not explain the free chain orientation. Another model based on anisotropic screening of excluded volume interactions has been proposed [21, 37]. To emphasise the role of orientational interactions in the appearance of the doublet in 2 H NMR spectra, model PDMS networks have been swollen to various extents with a good solvent of PDMS [53]. Figure 15.10 shows the slopes P=Δ/(λ2-λ-1) obtained in a PDMS network swollen with chloroform. The swelling process results in stretching

Figure 15.10 Slopes P=Δ/(λ2-λ-1) obtained on deuterated PDMS networks (Mn=10500 g.mol-1), swollen with chloroform, versus the polymer network volume fraction Φ (the slopes are normalised to the slope at Φ=1). Data from two networks with different characteristics are shown Reproduced with permission from P. Sotta, B. Deloche and J. Herz, Polymer, 1988, 29, 1171, Figure 2b. Copyright 1988, Elsevier Science

574

Deuterium NMR in Rubbery Materials network chains, and thus in enhancing the local average degree of orientation. This would lead to an increase of P, which is not observed. Instead, a decrease of the slope P is observed. This has been interpreted within the mean-field model mentioned in Section 15.3.5, as the effect of a weakening of interchain interactions, due to the presence of solvent molecules which decreases the local density of chain segments [53]. There is another interesting illustration of the mean-field effect presented previously. End-linked PDMS networks with a bimodal distribution of precursor chains (Mn = 25000 and 3000 g.mol-1) with various fractions ΦL of long chains, have been prepared. Short or long chains may be observed separately, due to selective deuteration. It is remarkable that short and long chains exhibit the same doublet in a given network [54].This shows that the doublet originates in a collective effect, which affects both type of chains in the same way, with the same magnitude. The doublet splitting obeys an ‘ideal mixing law’: Δ=ΦLΔL+(1-ΦL)ΔS; it is an average of that ΔL observed in the network with 100% long chains, and that ΔS in the network with 100% short chains.

15.3.7 Correlation to Elastic Properties In this section it is shown that 2H NMR is an excellent and quite direct way to characterise elastic properties of a rubber material. It has been mentioned in Section 15.3.5, that this characterisation may be performed in two ways. First, analysing the line shape, (i.e., measuring the linewidth), in the relaxed state gives the magnitude of local constraints stored in elastic chains, and this in turn is directly related to macroscopic elastic properties of networks. Alternatively, the slope P=Δ/(λ2-λ-1) is a function of parameters which are involved in elasticity as well. Both quantities contain the dominant factor 1/N, where N is the average length of elastic chains, and they should be correlated in a series of samples with various characteristics. This has been checked. However, due to the presence of dangling chains and perhaps other defects such as loops, it was shown that the line shape has to be analysed very carefully to establish such a correlation [25]. Indeed, the line shape allows the probing of the full distribution of local constraints (which may be quite inhomogeneous), whereas the doublet splitting only reflects a mean-field average in the system (Section 15.3.5). To probe the 1/N dependence in the slope P, experiments were carried out on model endlinked networks with different precursor chain lengths (average molecular weight Mn varying from 3100 to 23000 g.mol-1) [55, 56]. As expected, the slope P decreases when N increases. However, it does not follow a 1/N variation: P varies only by a factor of about 3 as N varies by a factor of 7. In an end-linked network, N should not be simply assimilated with the length of precursor chains, since trapped entanglements may play the role of effective crosslinks (often denoted as physical crosslinks), in addition to chemical

575

Spectroscopy of Rubbers and Rubbery Materials crosslinks [57]. Therefore, some effective mesh size Neff should be introduced to take the actual density of crosslinks into account [58]. A way to vary the number of trapped entanglements is to perform the crosslinking reaction at different polymer concentrations vc [14]. Indeed, network elastic properties depend on vc : the higher vc, the higher the elastic modulus and the polymer volume fraction at swelling equilibrium Φe [59, 60]. Amongst the various PDMS networks studied, those which differ by vc only do exhibit different slopes P [61, 62]. In Figure 15.11 the slope P is plotted versus Φe [57] for a series of networks with various crosslink and/or trapped entanglement densities. At equilibrium swelling in a good solvent, each chain of length N (between two adjacent junctions) occupies a volume of the order R3 = (aN3/5)3, which leads to Φe = N-4/5 [63] and thus P = Φe5/4. Though the dispersion of points in Figure 15.11 is too wide to test this power law dependence, there is a clear relationship between both quantities [25, 57]. On the other hand, Figure 15.12 shows the dependence of the slope P on the elastic modulus in a series of PDMS networks with various precursor chain lengths and elastic

Figure 15.11 Slopes P=Δ/(λ2-λ−1) as a function of the polymer volume fraction at swelling equilibrium in cyclohexane Φe in a series of end-linked PDMS networks with various lengths between junctions (A, A′: Mn = 25000 g.mol-1; B, B′: Mn = 10500 g.mol-1; C: Mn = 3100 g.mol-1) and polymer concentration during crosslinking in toluene (A, B, C: vc = 0.7; A’, B’: vc =1) Reproduced with permission from A. Dubault, B. Deloche and J. Herz, Progress in Colloid and Polymer Science, 1987, 75, 45, Figure 3. Copyright 1987, Spriger-Verlag GmbH & Co KG

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Figure 15.12 Slopes P=Δ/(λ2-λ-1) plotted versus the effective molecular weight between crosslinks Mc and the elastic modulus Ge, obtained from swelling measurements, in series of PDMS networks with various precursor chain lengths and elastic chain fractions: ● network chains only are deuterated (observed); O dangling chains only are deuterated; ❑ free probe chains only are deuterated Reproduced with permission from K. McLoughlin, J.K. Waldbieser, C. Cohen and T.M. Duncan, Macromolecules, 1997, 30, 1044, Figure 6. Copyright 1997, American Chemical Society

chain fractions [17]. There is a linear correlation between both quantities, as it is expected since both quantities are proportional to 1/Neff. In PB networks, the variation of the line shape as a function of the applied stress was interpreted in terms of a chain length distribution. Shorter chains may be more oriented than longer ones, at a given elongation [18], which may lead to a non-affine behaviour at the chain scale. The question of the spatial scale to which the deformation is affinely transmitted, has been investigated intensively by small angle neutron scattering [64]. However, it may happen as well that the chain portions close to junction points are more oriented (have a more restricted mobility) than those in the middle of the chains [19]. In Figure 15.13, the slopes P=Δ/(λ2-λ-1), measured in end-linked PDMS networks with bimodal distributions of precursor chain lengths and various types of crosslinking agents,

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Figure 15.13 Slopes P=Δ/(λ2-λ-1) versus the elastic modulus 2(C1+C2) (C1 and C2 are the Mooney-Rivlin coefficients), in bimodal PDMS networks composed of a mixture of long (Mn = 25000 g.mol-1) and short chains (Mn = 3000 g.mol-1). Either long (capital letters) or short (small letters) chains are deuterated. The long chain volume fractions are 15% (d, D), 50% (m, M) and 75% (s, S) Reproduced with permission from B. Chapellier, B. Deloche and R. Oeser, Journal de Physique II (Paris), 1993, 3, 1619, Figure 5. Copyright 1993, EDP Sciences

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Deuterium NMR in Rubbery Materials are plotted versus the elastic modulus (expressed in terms of the Mooney-Rivlin coefficients C1 and C2 as 2(C1+C2) [33]), which have been measured independently for the whole series of samples. The variation of the slope P is indeed perfectly linear in the elastic modulus, which corroborates very nicely the mean-field model presented in Section 15.3.5 [54]. The results reported in Figures 15.11, 15.12 and 15.13 show relationships between microscopic data (molecular orientation) and macroscopic ones (elastic modulus). This suggests that the orientational interactions introduced in Section 15.3.6 to account for the 2H NMR doublet induced under stress, play a role in rubber elasticity. Indeed, the force to apply to realise a given deformation is expected to be lower in the presence of interactions which tend to orient chain segments along the force direction. An elementary, phenomenological model of rubber elasticity including short-range interactions has been proposed [65]. It predicts the λ-dependence of the reduced force (and of the dilation modulus) and reproduces fairly accurately the elastic response of PDMS networks [66].

15.4 Using Deuterated Probes The necessity to label (deuterate) the systems under study may often appear as a disadvantage, since it prevents working with real (industrially designed) materials. A way to circumvent this drawback is to use labelled (deuterated) molecules as 2H NMR probes of material properties [67]. The uniaxial contribution to the stress-induced orientation (the second term in Equation 15.4) has been attributed mainly to orientational interactions between chain segments. When guest molecules are introduced in a rubber matrix, these interactions take place between network chain segments and guest molecules as well. This effect has been recognised earlier, and may be used to study indirectly the behaviour of the matrix. Two types of guest molecules have been used for this purpose. One possibility is to use free polymer chains, chemically identical to the network chains. In order to diffuse in the network easily, the free chains should not be too long, typically significantly smaller than the mesh size of the network under study. The advantage here is that guest molecules are locally identical to network chains and thus behave as ideal, nonperturbative NMR probes. Indeed, it was shown that free chain segment exhibit the same degree of uniaxial order as network chain segments, with the same strain dependence (see Figure 15.7) [39]. These properties remain valid when oligomers are used as probe molecules, which suggests that orientational couplings between chain segments are short range. The presence of probe chains, identical to network chains does not affect the local chain segment density. Thus, the induced order remains unaffected up to a quite large fraction of free chains or oligomers (about 25%) [53]. End-linked PDMS [38, 39], poly(urethane) (PU) networks [39], as well as PB networks [20] have been investigated in this way.

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Spectroscopy of Rubbers and Rubbery Materials Another possibility is to use deuterated molecules of a good solvent for the rubber matrix under investigation. The idea here is to introduce a relatively low fraction of solvent, so that the matrix properties are not affected significantly. The appearance of a well resolved doublet Δ in the NMR spectra of solvent molecules as the rubber is uniaxially stretched, shows that the solvent molecules are also sensitive to the uniaxial anisotropy induced in the deformed rubber [67]. This phenomenon is quite general and it has been observed with a number of polar or non-polar solvents (short alkane molecules, benzene, chloroform, cyclohexane, toluene, water). Indeed, it was shown that the induced orientational order, as measured in the doublet splitting, is extremely sensitive to the presence of even a small fraction of solvent: for example, in the system PDMS network/ chloroform, the orientational order induced on solvent molecules drops by about 20%, when the solvent volume fraction varies from 0% to 10%. The solvent degree of order is always smaller than the chain segment order [9]. However, due to the high mobility of solvent molecules, the linewidth observed in the solvent spectra is always very small, which allows to observe resolved doublets even with very small splittings (typically a few Hz) at small deformations (smaller than 5%, i.e., λ≤1.05) (Figure 15.14). Note also that, due to this mobility, only an average of the local orientation in the network is reflected by the solvent. In the low deformation limit, it has been shown that the solvent doublet is proportional to that of network chains (and thus exhibits the λ2-λ-1 dependence), with a proportionality constant which depends on the pair polymer/solvent only [9, 53]. Solvent molecules are thus very sensitive probes of the rubber matrix behaviour, and a number of studies dealing with network parameters such as the density of crosslinks or trapped entanglements have been achieved in this indirect approach [55, 61]. Anisotropic swelling of a rubber sample is another way to obtain uniaxial deformations. This may be achieved by swelling a cylindrical test sample, while maintaining its fixed diameter (sample confined in a tube). Such confined swelling experiments are convenient to deform uniaxially fragile systems as gels or physical networks. A deuterated solvent may be used as both the constraint agent and the 2H NMR probe allowing measurement of the induced uniaxial anisotropy. This experiment has been performed in various types of rubbers [68, 69]. It may be used to monitor the change in anisotropy as the constraint is released and its temporal relaxation as equilibrium is approached. The orientation of the swelling agent (solvent or free chains) has to be taken into account in the analysis of the stress-optical behaviour of swollen networks. Specifically, the segment polarisability (relative to the network chains or to the diluent chains), as currently derived from stress-optical coefficients [33], may not be representative of intrinsic properties of isolated chains. Short-range orientational interactions between the probe molecules and network chains (and between the chains of the matrix itself) must be considered in the interpretation of opticoelastic properties of swollen (and dry) rubbers [67].

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Figure 15.14 2H NMR spectra of deuterated benzene molecules diffusing in a uniaxially stretched PB network for various elongation ratios. The solvent volume fraction is 9% Reproduced with permission from A. Dubault, B. Deloche and J. Herz, Polymer 1984, 25, 1405, Figure 1. Copyright 1984, Elsevier Science

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15.5 Filled, Composite and Thermoplastic Elastomers 15.5.1 Filled Elastomers

15.5.1.1 Carbon Black Filled Vulcanised Natural Rubbers Deuterium has a quite low natural abundance (1.5 x 10-2 %). However, the main problem to obtaining measurable 2H NMR signals is related mainly to the rather broad linewidth in systems with interesting rubbery properties, (i.e., with enough local constraints stored in elastic chains, thus, large residual quadrupolar interactions), rather than to the small natural abundance itself. This can be circumvented by using the magic angle spinning (MAS) technique, which allows concentration of the widely spread spectral intensity into a family of sharp spinning sidebands. This has been used in a series of filled and unfilled sulfurvulcanised natural rubber (cis-polyisoprene) samples [70]. Figure 15.15 show a series of natural abundance 2H spectra obtained under MAS at a spinning speed of 0.5 kHz, in natural rubbers with various proportions of vulcanisers and carbon black fillers. The spectra are fitted with a two-component model which includes a ‘mobile’ (narrower) component and a ‘rigid’ (broader) one. The residual interactions obtained from best fits of the spectra varies from 0.2 to 0.6 kHz (mobile component) and from 1.2 to 1.7 kHz (rigid component). Thus, this experiment demonstrates that it is possible in principle to estimate the magnitude of local chain orientation using natural abundance 2H NMR.

15.5.1.2 Carbon Black Filled Diene Rubbers 1

H and 2H NMR have been used in styrene-butadiene rubber (SBR) with and without carbon-black fillers to estimate the values of some network parameters, namely the average network chain length N. The values obtained from both approaches were checked to make sure that they were consistent with each other and with the results of other methods [71, 72, 73]. To this purpose, a series of samples with various filler contents and/or crosslink densities were swollen with deuterated benzene. The slopes P=Δ/(λ2-λ-1) obtained on deuterated benzene in uniaxially stretched samples were measured. The slopes increase significantly with the filler content, which suggests that filler particles act as effective junction points [72, 73].

Deuterated PB networks filled with carbon black have been investigated recently [74]. The 2H NMR lineshape is different from that in unfilled elastomers: an asymmetric doublet is observed as the sample is uniaxially stretched (λ=1.8)). This asymmetry is related to the presence of carbon black fillers, which induce magnetic inhomogeneities.

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Figure 15.15 Natural abundance 2H MAS spectra observed in a series of vulcanised natural rubbers with various vulcaniser (sulfur: 1 or 3 wt%) and/or filler (carbon black: 0 or 40 wt%) contents. The spinning speed is 0.5 kHz. The number of scans is about 300000. Spectra are simulated with two components (a ‘mobile’ and a ‘rigid’ one) with various residual quadrupolar interactions Reproduced with permission from C. Malveau, P. Tekely and D. Canet, Solid State Nuclear Magnetic Resonance, 1997, 7, 271, Figure 5. Copyright 1997, Elsevier Science

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Spectroscopy of Rubbers and Rubbery Materials This effect allows to quantify the fraction of polymer in close contact with filler particles. It is strongly suggested that that fraction is submitted to local deformations different from that within the bulk matrix.

15.5.1.3 Silica Filled Silicon Rubbers In composite systems, 2H NMR is particularly suited to investigate interfacial properties. Indeed, isolated nuclei are observed, which potentially allows spatially selective information to be obtained. It has been used to investigate polymer chain mobility at the polymer-filler interface, mainly in filled silicon (in particular PDMS) networks. The chain mobility differs considerably at the polymer-filler interface, and this may be interpreted in terms of an adsorbed polymer layer at the filler surface. T1 relaxation measurements allowed to determine the fraction of chain units involved in the adsorption layer, or equivalently, the thickness of the layer [75, 76, 77]. The molecular mobility and the thickness of the adsorption layer are very sensitive to the type of filler surface [78]. It was shown that the stress-induced orientational order is larger in a filled network than in an unfilled one [78]. Two effects explain this observation: first, adsorption of network chains on filler particles leads to an increase of the effective crosslink density, and secondly, the microscopic deformation ratio differs from the macroscopic one, since part of the volume is occupied by solid filler particles. An important question for understanding the elastic properties of filled elastomeric systems, is to know to what extent the adsorption layer is affected by an external stress. Long-time elastic relaxation and/or non-linearity in the elastic behaviour (Mullins effect, Payne effect) may be related to this question [79]. Just above the melting temperature Tm, it has been shown that local chain mobility in the adsorption layer decreases under stress, which may allow some elastic energy to be dissipated, (i.e., to relax). This may provide a mechanism for the reinforcement of filled PDMS networks [78]. Silica hydrolysis-condensation of alkoxysilane-terminated chains of various polymers (PB, poly(propylene)-glycol, poly(ethylene)-glycol) leads to the formation of silica nodules. Hybrid (organic-inorganic) nanocomposite rubbery materials are obtained in this way. These materials have been swollen by various deuterated solvents, in order to probe the order induced in the polymer matrix under stress [80].

15.5.2 Semi-crystalline Polymers In semi-crystalline polymers, crystallites may act as effective crosslinks, which gives rise to rubbery properties above Tg. In systems with long chains, the presence of crystallites 584

Deuterium NMR in Rubbery Materials allows to relax the constraints in the amorphous phase, while the chains may be quite strongly oriented as regards their overall configuration. Deuterated n-octane was used as a probe of the amorphous regions in a stretched PDMS network cooled down below Tm (octane molecules do not penetrate into crystallites). It was observed that the local chain orientation in the amorphous regions decreases as the system crystallises [81]. Drawn isotactic polypropylene (iPP) fibres have been studied by using deuterated ndecane as a 2H NMR probe of the chain deformation ratio in the amorphous regions. It is observed that the slope P=Δ/(λ2-λ-1) (defined in the limit of low deformations) indeed depends on the annealing temperature [82]. Thus, annealing above the melting temperature Tm of the crystallites allows chains to relax to some extent. Then, the local deformation ratio in the amorphous phase λl becomes lower than the macroscopic one λ and depends on the annealing temperature, i.e., on the amount of chain relaxation. Therefore, such systems have strongly non-affine deformation at the chain scale. Fully deuterated linear poly(ethylene) (PE) has been investigated also via various 2H NMR techniques below Tm, i.e. in the semi crystalline state [83, 84]. The crystallinity ratio was measured as a function of temperature and it was shown that the motions are highly restricted in the amorphous regions of PE. It was shown that the onset of β and α transitions (at which mobility appears in the crystalline phase) may be observed by 2H NMR on raising the temperature. This onset of local motions in the crystalline phase is related to the chain relaxation process quoted previously.

15.5.3 Thermoplastic Elastomers Block copolymers allow to realise thermoplastic elastomers, by combining a central, soft block with rigid (crystalline and/or glassy) blocks at chain extremities. The rigid segments segregate into hard microdomains (or nodules) which act both as crosslink points with high functionality and as solid, reinforcing particles.

15.5.3.1 Polyamide–Polyether Block Copolymers Polyamide–polyether (PA-PEG) block copolymers have been studied using deuterated water (2H2O) molecules as 2H NMR probes [85, 86, 87]. Well resolved doublets are observed on 2H2O on stretching the samples at room temperature. At a given elongation ratio λ, the splitting depends on the water concentration c. The linewidth is that of a liquid and becomes comparable to that in bulk water for c≈0.6. This confirms that water

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Spectroscopy of Rubbers and Rubbery Materials probes the mobile, elastomeric part of the sample only (PEG block). Moreover, the linewidth follows the Williams-Landel-Ferry (WLF) equation and is independent of the stretching ratio and the size of the hard PA block. In these systems, the 2H 2O splitting is not linear in the strain function (λ2-λ -1) (Figure 15.16). This indicates that the microscopic deformation ratio λl (at the chain

Figure 15.16 2H NMR splittings in deuterated water molecules in a PA-PEG blockcopolymer versus the stretching functions λ2-λ-1 (a) or λ*2-λ*-1 (b) (see Section 15.5.3.1), at various water concentrations c (the points at c = 0% are deduced from the measured splittings Δ(c) by extrapolation) Adapted with permission from J. Rault, C. Macé, P. Judeinstein and J. Courtieu, Journal of Macromolecular Science B, 1996, 35, 115, Figure 6. Copyright 1996, Marcel Dekker, Inc

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Deuterium NMR in Rubbery Materials scale) is different from the macroscopic one λ, or in other words, that the deformation is not affine at the chain scale. This is a quite general property in thermoplastic elastomers. As PA-PEG elastomers are stretched (at room temperature) and then annealed just above Tm, they do not relax to their initial dimension. This indicates that flow occurs during deformation, because PA segments recrystallise with the chain axis parallel to the stretch direction. However, the splitting variation tends to be linear in (λ*2-λ*-1), in which λ* is the deformation ratio referred to the relaxed dimension measured after annealing (Figure 15.16) [87]. This indicates that the local deformation ratio λl is equal to λ*. Note additionally that 2H2O has be used as well to probe PA-PEG block-copolymers doped with several caesium salts [88].

15.5.3.2 Polyurethane Elastomers PU elastomers based on a polyester diol as the soft block and isocyanate groups coupled via a chain extender as the rigid block, have been investigated. The role of the chain extender is to induce some mobility and disorder within the hard microdomains. The elastomers are selectively deuterated in the rigid block. 2H NMR spectra show the coexistence of a solid and a mobile fraction [89]. The solid fraction shows no onset of fast molecular motions at all, while the mobile fraction increases with temperature. These results suggest that some of the hard microdomains (perhaps the most disordered ones) melt successively when temperature is raised. PU systems with hard blocks made of piperazine, have been also studied [90, 91, 92]. In systems deuterated in the segments of the soft blocks, 2H NMR results show that a fraction of the soft segments have a restricted mobility, due to the connectivity to hard blocks, rather than to the adsorption on those blocks acting as filler particles [93]. To study the interface in a specific way, systems with hard blocks selectively deuterated at different positions have been synthetised. A reduction of quadrupolar interaction according to the position was observed, corresponding to an increase of the mobility in hard segments on approaching the hard-soft interface.

15.5.3.3 H-bonded Polybutadiene Networks New crosslinked elastomers may be formed by self-aggregation of substituted functional units able to form directed hydrogen-bonds. 1,4-PB has been randomly substituted with about 4 mol% urazole units [94]. The urazole units form hydrogen-bonded supramolecular, plate-like aggregates which act as effective crosslinking zones, thus creating a thermoplastic elastomer [95, 96, 97]. The system is deuterated either in the PB backbone or in the urazole units, which allows to investigate the molecular dynamics in

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Spectroscopy of Rubbers and Rubbery Materials either part of the elastomer. The 2H NMR spectra in the relaxed state show motional heterogeneities along the PB chains as well as in the junction domains. About 97% of the urazole units are included in the aggregates, in which their mobility is completely frozen. There are however some phenyl ring flips at the boundary of the domains. The 3% urazole units which are not included in the aggregates undergo fast isotropic motions. Part of the PB segments in the elastic chains are restricted in their mobility: a layer of about two monomers units adjacent to the aggregates is immobilised. The stress-induced quadrupolar doublets have been observed in samples deuterated in PB chains [98]. At high elongation ratios, the doublet splittings are not proportional to the strain function (λ2-λ-1): as λ increases, the splittings first increase rapidly, then become much less dependent on λ or even nearly constant in some cases. There is also a huge increase in linewidth on increasing λ. These features have been tentatively interpreted as a very broad distribution of chain segment orientation in the highly stretched samples [96].

15.6 Conclusion The results discussed in this chapter demonstrate that 2H NMR is a powerful technique for investigating microscopic properties in rubber networks. Most of the experiments described here are easy to handle on standard NMR equipment. Due to the absence of interactions between 2H nuclei, spectra and line shapes are easy to interpret and give quite direct information, at least in the first step of analysis, which is that generally required to correlate microscopic to macroscopic properties in these systems. Additionally, in contrast to optical techniques (as birefringence, infrared dichroism, fluorescence polarisation) the information which is obtained is very specific, because spatial and temporal averaging processes are clearly distinguishable in NMR. These basic properties have been exploited in the experiments described here. Various time scales have been investigated. The static or dynamic nature of the observed local anisotropy has been specified. It has been demonstrated that 2H NMR is sensitive to very low degrees or very small variations in the magnitude of the elastic constraints stored in elastic chains. The reliability and sensitivity of the probe method has been emphasised. It circumvents almost completely the perturbations inherent to some other probe techniques (electron spin resonance, fluorescence). In particular, free chains appear to be ideal, non-perturbative NMR probes for testing chain segment orientation in strained rubbers. The solvent probe method is easy to handle and unexpensive, since it does not require the synthesis of

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Deuterium NMR in Rubbery Materials labelled polymers. In view of industrial applications, the technique may be used in nontransparent samples, such as multi-component or filled elastomers. This technique should be widely used to characterise rubbery materials, as regards for example the presence of residual constraints or deformations resulting from material processing and/or from intensive mechanical deformations. One of the most relevant physical results discussed here is that the stress-induced orientation in rubbers is not a single chain process, but involves cooperative interactions between chain segments. Indeed, the 2H NMR response of elastic chains in a uniaxially strained rubber is similar to that of a uniaxial fluid. This basic result is in contrast to various theoretical descriptions wherein the physical effects of the chains are mainly restricted to their sole action on the junctions. These local interactions between chains may affect macroscopic elastomer properties (elasticity, optoelasticity, resilience) and thus must be taken into account in order to interpret these properties at the microscopic level. For example, a way to change (and perhaps to enhance) drastically the properties of rubbers would be to enhance the orientational interactions by introducing liquid crystalline (nematogenic) groups along the chains [99, 100]. The collective, dynamical behaviour of chain segments observed in strained rubbers is a general property of molten polymer systems in which the chains are constrained or confined, such as liquid crystal polymers [101], block copolymers [102], grafted polymer brushes [10,103], coated polymer films [104] and melts under shear flow [105]. More specific 2H NMR investigations, perhaps using selective deuteration of network chains, may allow to obtain a very precise description of rubbery materials at a molecular scale, as regards the characterisation of the heterogeneity in the local order or the dynamics. The distribution of local order (or of the local degree of constraint) is probably related to the heterogeneity of the crosslink density, or equivalently to the distribution of elastic chain lengths. Some degree of motional heterogeneity may be associated to the presence of junctions in model (single component) networks. On the other hand, a strong gradient of chain segment mobility is related to the presence of polymer-solid interfaces in filled and/or thermoplastic (multicomponent) elastomers. The mechanical properties of these rubbery materials depend in a crucial way on both properties, namely, the repartition of (effective) crosslinks and/or the nature of the interfaces in filled systems. Note finally that, due to its high specificity and sensitivity to the local degree of constraints in a rubbery material, 2H NMR has been used to obtain a spatial map of the constraints in a material by imaging techniques [106, 107]. These experiments are described in another chapter of this book. 2H NMR has been used as well to investigate other related polymer systems, such as gels [108], thermosets [109], fibres [110, 111, 112, 113] and various solid (glassy and/or semi crystalline) polymers [114].

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595

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596

Abbreviations

Abbreviations

1D

Two-dimensional

1Q

Single-quantum

2D

One-dimensional

2

D-NMR

Deuteron solid-state NMR

2

H 2O

Deuterated water

2Q

Double-quantum

3Q

Triple-quantum

9BEN

9 borobicyclo [3, 3, 1] nonane

A1

Amide 1

ABS

Acrylonitrile-butadiene-styrene terpolymer

AC

Accelerator

ACM

Acrylate rubber

ACN

Acrylonitrile

AFM

Atomic Force Microscopy

Al-CSM

Aluminium salt of chlorosulfonate polyethylene

ALMA-d5

Allyl-d5 methacrylate

Al-m-EPDM

Aluminium salt of maleated EPDM

AN

Acrylonitrile

AO

1-Allyl oxyoctane

APT

Attached Proton Test

ASTM

American Standards for Testing and Materials

ATR

Attenuated total reflectance spectroscopy

ATR-IR

Attenuated total reflectance-IR spectroscopy

B/S

Butadiene/styrene

BABA

Back-to-back

597

Spectroscopy of Rubbers and Rubbery Materials BDO

1, 3-Butadiene

BET

Brunauer, Emmett and Teller

BHT

Butylated hydroxytoluene

BPO

Benzoyl peroxide

BR

Polybutadiene rubber

Bt

Benzothiazole

CB

Carbon black

CBR

Cyclised polybutadiene rubber

CBS

N-cyclohexyl-2-benzothiazole sulfenamide

CC

Crosslinks

CHD

Cis-1,4 hexadiene

CIIR

Chlorobutyl rubber

CK

Co-operative kinematics

CM

Chlorinated polyethylene

CNR

Chlorinated natural rubber

COLOC

Chemical shift correlation

COSY

Correlation Spectroscopy

CP

Cross Polarisation

CPMG

Carr, Purcell, Meiboom and Gill

CR

Chloroprene rubber

CRAMPS

Combined Rotation and Multiple Pulse Spectroscopy

CSDPF

Carbon-silica dual phase fillers

CSM

Chlorosulfonated PE

CT

Cold trap

DAIP

Diallylisophthalate

DANTE

Delays alternating with nutations for tailored excitations

DAO

1,8-diallyl oxyoctane

DAP

Diallylphthalate

DATP

Diallylterephthalate

DBM

Dibutyl maleate

DBTDL

Bibutyl tin dilaurate

DBTM

Dibutyl tin maleate

DCDPH

Selectively hydrogenated DCPD

DCE

Dipolar correlation effect

598

Abbreviations

DCP

Dicumyl peroxide

DCPD

Dicyclopentadiene

DD

Dipolar decoupling

DEPT

Distortionless enhancement by polarisation transfer

DETA

Dielectric thermal analysis

DIN

Deutsches Institut für Normung

DMA

Dynamic mechanical analysis

DMS

Dimethyl siloxane

DOS

Disk operating system

DPNR

Deproteinised NR

DQ

Double quantum

DRIFTS

Diffuse reflectance IR-FT Spectroscopy

DSC

Dynamic scanning calorimetry

DTA

Differential thermal analysis

DTG

Differential thermal gravimetry

DTGS

Deuterated triglycine sulfate

EAA

Polyethylene-co-acrylic acid

EBA

Polyethylene-co-butyl acrylate

ECO

Epichlorohydrin

EGP

Evolved gas profile

EHA

2-Ethyl hexyl acrylate

EN

Chain entanglement

EN/TE

Temporary chain entanglements

EN/TR

Trapped chain entanglements

ENB

5-Ethylidene-2-norbornene

ENBH

2-Ethylidene norbonane

ENR

Epoxidised natural rubber

ENVP

EPDM-g-NVP

EP

Ethylene-propylene copolymer(s)

EPDM

Ethylene-propylene diene terpolymer(s)

EPM

Ethylene-propylene copolymer(s)

EPR

Ethylene-propylene rubber

ER

Polyurethane-epoxide resin

ESBS

Epoxidised butadiene-styrene block copolymer

599

Spectroscopy of Rubbers and Rubbery Materials ESO

Epoxidised soya bean oil

EV

Efficient vulcanisation

EVA

Ethylene-vinyl acetate copolymer

EVS

Polyethylene-co-vinyl trimethoxy silane

FDP

Farnesyl diphosphate

FGP

Functional gas profiles

FID

Flame ionisation detection

FIND

Free induction decay(s)

FRLS

Fire-resistant low smoke

FT

Fourier transform

FT-IR

Fourier transform-infrared spectroscopy

GC

Gas chromatography

GHPD

Gated high power decoupling

GMA

Glycidyl methacrylate

GPC

Gel permeation chromatography

HBR

Hydrogenated butadiene rubber

HCR

Hydrogenated CR

HD

1,4-hexadiene

HEMA

2-hydroxy ethyl methacrylate

HENR

Hydrogenated epoxidised natural rubber

HETCOR

Heteronuclear chemical shift correlation

HMQC

Heteronuclaer multiple-quantum coherence

HNBR

Hydrogenated nitrile rubber

HNR

Hydrogenated natural rubber

HOFR

Heat, oil and fire resistant

HPLC

High pressure liquid chromatography

HPV

High pressure vulcanisation

HQEE

Hydroquinone bis (b-hydroxyethyl) ether

HRTG-EI/SI QMSHigh resolution thermogravimetry–electron impact/soft ionisation quadrupole mass spectrometry HSBR

Hydrogenated styrene butadiene rubber

HTNR

Hydroxyterminated liquid NR

HXSBR

Hydrogenated carboxylated styrene butadiene rubber

ICTA

International Confederation for Thermal Analysis

600

Abbreviations

IDP

Isopentenyl diphosphate

IIR

Polyisoprene rubber

IIRP

International Institute of Synthetic Rubber Producers

IPN

Interpenetrating network(s)

IPNF

Interpenetrating network foam

iPP

Isotactic polypropylene

IR

Infra red

ISAF

Intermediate super abrasion furnace

ITPE

Ionic thermoplastic elastomer

LDPE

Low density linear polyethylene

LDPE-g-PNVP

LDPE grafted with 2-N-vinyl pyrrolidone

LPA

Copolymer of polyamide

MA

Maleic anhydride

MAS

Magic angle spinning

MBS

2(4-Morpholinothio) benzathiazole

MBT

Mercaptobenzothiazole

MDI

Methylene bis-4-phenol isocyanate

MDPE

Medium density polyethylene

MDSC

Modulated differential scanning calorimetry

MEPDM

Maleated EPDM

MMA

Methyl methacrylate

Mn

Molecular weight

MNB

5-Methylidene-2-norbonene

MOR

N-oxydiethylene-2-benzothiazole sulfenamide

MP

Melting point

MPTES

Mercaptopropyl triethoxysilane

MQ

Multiple quantum

MRI

Magnetic resonance imaging

MS

Mass spectrometry

MSA

Maleic anhydride grafted polypropylene

m-SEBS

Maleated styrene-ethylene-butylene-styrene block copolymer

MS-PCA

Mass spectrometry with principal component analysis

MW

Molecular weight

Na-m-SEBS

Sodium salt of maleated styrene ethylene butylene styrene

601

Spectroscopy of Rubbers and Rubbery Materials Na-s-m-SEBS

Sodium salt of sulfonated maleated SEBS

NBR

Nitrile-butadiene rubber

NIST

National Institute of Science and Technology

NMA

N-methylol acrylamide

NMR

Nuclear magnetic resonance

NMR-MOUSE

NMR-mobile universal surface explorer

NOE

Nuclear Overhauser effect

NOESY

NOE spectroscopy

NR

Natural rubber(s)

NVP

N-vinyl pyrrolidine

OIT

Oxidative induction time

OPD

Optical path difference

PA

Photoacoustic

PA-12

Polyamide-12

PA-FTIR

Photoacoustic-Fourier transform-Infrared spectroscopy

PA-PEG

Polyamide–polyethylene glycol

PAS

Photoacoustic spectroscopy

PB

Poly(1,4-butadiene)

PBT

Polybutylene terephthalate

PC

Polycarbonate

PCA

Principal component analysis

PCB

Polychlorinated biphenyl(s)

PCL

Polycaprolactone

PCP

Polchloroprene

PDMS

Polydimethylsiloxane

PDSC

Pressure differential scanning calorimetry

PDXL

Polydioxolane

PE

Polyethylene

PED

Deuterated PE

PEEK

Polyether ketone

PEG

Polyethylene glycol

PEGDA

Poly(ethylene glycol)diacrylate

PEgDBM

PE grafted to dibutyl maleate

PEH

Normal PE

602

Abbreviations

PEO

Polyethylene oxide

PEXa

Peroxide crosslinking system

PEXb

Silane crosslinking system

PEXc

Electron beam crosslinking system

PF

Phenolic formaldehyde

PFG

Pulsed field gradients

PFGE

Pulsed field gradient echo spectroscopy

PFPA

Pentafluorophenyl acrylate

PFPMA

Pentafluoro phenyl methacrylate

phr

Parts per hundred rubber

PIB

Polyisobutylene

PMDETA

Penta-methyl diethylene triamine

PMMA

Polymethylmethacrylate

PP

Polyopropylene

PPA

Polypentenamer

PPL

Poly-b-propiolactone

ppm

Parts per million

PPO

Poly(propylene oxide)

PR

Alkyl phenol formaldehyde resin

PS

Polystyrene

PTA

Pulsed thermal analysis

PTFE

Polytetrafluoroethylene

PTMG

Poly(tetramethylene ether glycol)

PTMO

Polytetramethyleneoxide

PTOP

Para-tert-octyl phenolic

PU

Polyurethane(s)

PUA

Polyurethane acrylate

PUF

Polyurethane foam

PVC

Polyvinyl chloride

PVME

Polyvinyl methyl ether

PyGC-FID

pyrolysis gas chromatography – FID

QMS

Quadrupole mass spectrometry

R

Dichroic ratio

rf

Radiofrequency

603

Spectroscopy of Rubbers and Rubbery Materials RG

Rosencwaig and Gersho theory

RX

Receiver signal

S/N

Signal-to-noise ratio

SAN

Styrene acrylonitrile

SBR

Styrene-butadiene rubber(s)

SBS

Styrene-butadiene-styrene rubber

SEBS

Styrene-ethylene-butylene-styrene block co-polymer

SEC

Size exclusion chromatography

SEM

Scanning electron microscopy

SEV

Semi-efficient vulcanisation

SGP

Specific gas profiles

SI QMS

Soft ionisation quadrupole mass spectrometry

SIS

Styrene-butadiene-styrene

SPRITE

Single-point ramped imaging with T1 enhancement

SQ

Single quantum

SRF

Semi-reinforcing furnace carbon black

STA

Simultaneous thermal analysis

TAC

Triallylcyanurate

TBBS

N-tertiary butyl-2 benzothiazole sulfenamide

TBSI

N-tertiary butyl-2 benzothiazole sulfenimide

TBTCl

Tributyltin chloride

TBTLi

Tributyltin lithium

TD

Thermal desorption

TDI

Toluene diisocyanate

TeCEA

Tetrachloroethyl acrylate

TESBT

Triethoxysilyl benzothiazole

TESPT

Bis(triethoxysilyl propyl) tetrasulfide

Tg

Glass transition temperature

TG

Thermogravimetry

TGA

Thermogravimetric analysis

THF

Tetrahydrofuran

Tm

Melting temperature

TMA

Thermomechanical analysis

TMBA

Thermo molecular beam analysis

604

Abbreviations

TMPS

Poly(tetramethyl-p-silphenylene-siloxane)

TMPTA

Trimethylol propane triacrylate

TMPTMA

Trimethylol propane trimethacrylate

TMTD

Tetramethyl thiuram disulphide

TOCSY

Total correlation spectroscopy

ToFMS

Time-of-flight MS

TPE

Thermoplastic elastomer(s)

TPEE

Thermoplastic polyester elastomers

TPGDA

Tripropylene glycol diacrylate

TPP

Triphenyl phosphate

TPU

Thermoplastic polyurethane

TRIM

Trimethylolpropanetrimethacrylate

TVA

Thermal volatilisation analysis

TX

Transmitter signal

U4A

4-(3,5-Dioxo-1,2,4-triazolidin-4-yl) benzoic acid

UV

Ultraviolet

VAC

Vinyl acetate

VI

Video imaging

VNB

5-Vinyl-2-norbornene triallylcyanurate

VTEO

Vinyl triethyl silane

VTMO

Vinyl trimethyl silane

WISE

Wide line separation

WLF

Williams-Landel-Ferry

Xe

Xenon

XNBR

Carboxylated nitrile rubber

XPS

X-ray photo-electron spectroscopy

XRD

X-ray diffraction

XSBR

Carboxylated styrene butadiene rubber

XSBR

Hydogenated, carboxylated SBR

ZDMC

Zinc dimethyl dithiocarbamate

Zn-EMA

Zinc salt of ethylene-methacrylic acid copolymer

Zn-m-EPDM

Zinc salt of maleated ethylene-propylene diene monomer rubber

Zn-m-HDPE

Zinc salt of maleated high density polyethylene

Zn-PEA

Zinc salt of polyethylene-co-acrylic acid

605

Spectroscopy of Rubbers and Rubbery Materials Zn-PPA

Zinc salt of polypropylene-co-acrylic acid

Zn-s-EPDM

Zinc salt of sulfonated EPDM

Zn-s-m-EPDM

Zinc salts of sulfonated maleated EPDM

Zn-XNBR

Zinc salt of carboxylated nitrile rubber

606

Contributors

Jean Pierre Cohen Addad Université Joseph Fourier - Grenoble, Laboratoire de Spectrometrie Physique, BP 87, 38402 St Martin d’Heres Cedex, Grenoble I, France

Jan C.J. Bart DSM Research, PO Box 18, 6160 MD Geleen, The Netherlands

Bernhard Blumich Institut für Technische Chemie und Makromolekulare Chemie, RWTH, Worringerweg 1, D-52056, Aachen, Germany

Prajna P. De Rubber Technology Centre, Indian Institute of Technology, Kharagpur 721302, West Bengal, India

Bertrand Deloche Laboratoire de Physique des Solides (CNRS UMR 8502), Université de Paris-Sud, 91405 Orsay Cedex, France

Dan E. Demco Institut für Technische Chemie und Makromolekulare Chemie, RWTH Aachen, Worringerweg 1, D-52056, Aachen, Germany

Herman G. Dikland DSM Elastomers, Global R&D, PO Box 1130, 6160 BC Geleen, The Netherlands

Martin van Duin DSM Research, PO Box 18, 6160 MD Geleen, The Netherlands

Eng Aik Hwee Ansell Shah Alam Sdn Bhd, Lot 16, Persiaran Perusahaan, Seksyen 23, 40000 Shah Alam, Selangor Darul Ehsan, Malaysia

607


Siegfried Hafner Varian Deutschland GmbH, Alsfelder Strasse, 64289 Darmstadt, Germany

Jack L. Koenig Case Western Reserve University, Department of Macromolecular Science, Kent Hale Smith Building – Room 212, 10900 Euclid Avenue, Cleveland, OH 44106-7202, USA

Victor M. Litvinov DSM Research BV, PO Box 18, 6160 MD Geleen, The Netherlands

Dallas D. Parker Case Western Reserve University, Department of Macromolecular Science, Kent Hale Smith Building – Room 212, 10900 Euclid Avenue, Cleveland, OH 44106-7202, USA

James R Parker PPG Industries Inc., Chemicals Group - Technical Centre, 440 College Park Drive, Monroeville, PA 15146, USA

Charles Raemaekers DSM Research, PO Box 18, 6160 MD Geleen, The Netherlands

Jitladda Sakdapipanich Department of Chemistry, Faculty of Science, Mahidol University, Rama VI Road, Bangkok 10400, Thailand

Nikhil K. Singha Department of Chemical Engineering, Division of Polymer Chemistry, Eindhoven University of Technology, PO Box 513, 5612 MB Eindhoven, The Netherlands

Paul Sotta Laboratoire de Physique des Solides (CNRS UMR 8502), Université Paris-Sud, 91405 Orsay Cedex, France

Hans W. Spiess Max-Planck Institut für Polymerforschung, Postfach 3148, D - 55021 Mainz, Germany

608

Contributors

Yasuyuki Tanaka Department of Chemistry, Faculty of Science, Mahidol University, Rama VI Road, Bangkok 10400, Thailand

Wiebren S. Veeman Institute of Chemistry, Gerhard-Mercator-University, Duisburg, Germany

Andrew K. Whittaker Centre for Magnetic Resonance, University of Queensland, QLD 4072, Australia

609


610

Index A ABS (acrylonitrile–butadiene–styrene) 21– 22, 26 Absorptivity 82 Accelerators IR spectra 24–25 sulfur vulcanisation 215, 322–323, 328–329, 330–333, 336, 506–507 Acid-catalysed cyclisation 420–421 Acrylonitrile, determination in NBR 61– 62, 88–90 Acrylonitrile–butadiene–styrene (ABS), thermal degradation 21–22, 26 Activators, sulfur vulcanisation 215, 323 Additives 27 Adhesion 111–113 Adsorption filler surface 354, 371–372, 375–377, 378–379 mineral aggregates 293–294, 315–317 Aflas FA-150C rubber, thermal degradation 21–22 Ageing elastomers 264–265, 274–275, 348 NMR imaging 264–265, 274–275 see also photoageing; thermooxidative degradation Aggregated polymers 294 Alkyl phenol formaldehyde resin (PR), IR spectra 94–95 Alkylthiols, EPDM crosslinking 222–223 1-Allyloxyoctane (AO) 230–232 Aluminium, polyethylene laminates 178– 179

Anisotropic interactions 322, 519–522, 524, 529, 558, 559–560 Anti-Markovnikov rule 142–143 AO (1-allyloxyoctane) 230–232 APT (attached proton test) 329 ASTM standards D297-93 14 D3677 83 E1131 14 PA–FTIR 64 temperature calibration 13 ATR see attenuated total reflectance (ATR) spectroscopy Attached proton test (APT) 329 Attenuated total reflectance (ATR) spectroscopy 81, 104–108, 109, 216, 227–228

B Balata cyclisation 142 IR spectra 88 isomers 88, 403, 431 BDO (1,4-butanediol polyurethanes) 189–192 Beer–Lambert’s law 82 Bell, Alexander Graham 49 Bending vibrations 79, 170–172 BHT (butylated hydroxytoluene) 237 Biocompatibility 557 Biosynthesis 402, 403 Blends 13 C NMR 340

611

Spectroscopy of Rubbers and Rubbery Materials crosslink density 502–510 1 H NMR transverse magnetisation relaxation 384 interfaces 271, 354 ionomeric 151–153, 154–155 IR spectra 92–98, 111–113, 180, 182– 183 mixing process 264 NMR imaging 264, 271, 272, 338, 384 orientation effects 71 polypropylene EPDM 482–485 EPM 465–474, 482 PVC 180, 182–187 self-crosslinking 96–98, 183, 186 thermoplastic elastomers 188, 189 Block copolymers EPM–polypropylene 465–474 IR spectra 98, 100 network structure 293 NMR 416, 419 PBT/PTMO 486–487 polyamide–polyether 585–587 see also copolymers Bloom (surface layer) 66–68 Bohr equation 77 Bond strength 80 BR see butadiene rubber Branch-points 168–169, 426–429, 439 Bromination, IR spectra 137 Butadiene rubber (BR) chain segment diffusion 499 crosslinking 227, 335–336, 337, 339 crosslink density 335–336, 506–507 cyclohexane self-diffusion 497–498 IR spectra 126, 127 chlorination 135–138 cyclisation 140–142 epoxidation 146 hydrocarboxylation 146

612

hydrogenation 129–131 hydrosilylation 142–144 isomers 83, 88, 90, 139, 212–213 NMR imaging 500–501 peroxide-curing 227, 338–339 sulfur vulcanisation 212–213, 333–338 swollen 500–501 1,4-butanediol (BDO) polyurethanes 189–192 Butylated hydroxytoluene (BHT) 237 N-t-butyl-2-benzothiazole sulfenamide (TBBS) 330–333, 336–337 Butyl rubber branch-point 426–427 1 H NMR 426–427 IR spectra 94–95, 127 unsaturation 405, 406

C Cable sheathing compounds 172, 175 Calibration IR spectroscopy 82 thermogravimetric analysis 13–14 Carbon black IR spectra 198–199 properties 465 Carbon-black-filled rubbers 81 ATR–FTIR spectra 105–106, 107 chain adsorption 317, 371–372 13 C NMR 341–347 chlorosulfonated polyethylene 58–60 crosslink density 507 diene rubbers 582, 584 DRIFTS 104, 105 EPDM 369–370, 372, 373, 463, 465, 474–480 free radicals 368, 369 1 H NMR transverse magnetisation relaxation 265–266, 369–374, 475 IR spectra 198–199

Index NBR 504 NMR-MOUSE 278–280 NR 341–344, 582, 583 PA–FTIR spectra 57–60 polymer-filler interactions 474–475 SBR 24–25, 57–58, 372, 374 sulfur vulcanisation 336–337, 341– 347, 582, 583 Carbon–carbon unsaturation 126, 128– 129 Carbon–silica dual phase fillers (CSDPF) 108 Carboxylated nitrile rubber (XNBR) IR spectra 85, 96–97, 99, 194–198 self-crosslinking blends 96–98, 99 Carboxylated styrene–butadiene rubber (XSBR), IR spectra 131, 133 Carboxylate ionomers, FTIR spectra 147– 149 Carboxynitroso rubber, thermal decomposition 28 Chain dynamics 312–315, 525 1 H NMR transverse magnetisation relaxation 313, 355–360, 361, 371, 374–376, 379 2 H NMR 562–566 polymer solutions 494–496 segmental mobility 495–496, 498– 499, 511–512, 520, 524–525, 530, 538, 548–549 time scale of motions 563–566 Chain structure adsorption 293–294, 315–317, 354, 371–372, 375–377, 378–379 branch-points 168–169, 426–429, 439 conformation 504 crosslinking 302–306, 493 dangling chains 563, 564, 571 disentanglement 361 elongation 296–298, 567–570, 573 end-groups 404, 405, 422–429, 436– 437, 438

end-linked networks 299, 561, 563, 575, 577–579 entanglements 309–315, 354, 360, 366–367, 377–378, 505–506, 561 extension 557–558 free chains 570–571, 579 linkage effect 302 long-chain dynamics 313–315 molecular mass distribution 364–366 network chain density 333, 334, 344, 346–347, 355–360, 367 polymeric networks 293–307 scission 424–426 short-chain dynamics 312 strand length dependence 301–302 swelling effects 302, 303, 304, 306, 344, 346–347 telechelic chains 299, 423–424, 439 terminal groups 424–426, 436–437, 438 thermal behaviour 301–302 see also molecular structure Characterisation chemical characterisation 321–400 chemically modified rubbers 125–165 elastomers 1–47 rubbery materials 125–165 chemical and physical networks 353– 400 molecular behaviour 557–558 soft polymeric matter 291–320 structural 17 viscoelastic materials 519–555 vulcanisates 321–352 Chemical analysis chemical composition distribution 410–413 elastomers 402–422 (multi) hyphenated TGA 17 network structure 354–355 Chemical degradation 442 Chemically modified rubbers

613

Spectroscopy of Rubbers and Rubbery Materials IR spectroscopy characterisation 125– 165 NMR 419–422 Chemical networks see crosslinking Chemical shift 13 C NMR spectroscopy 325–326, 328, 524, 543 Fourier transform NMR 508 NMR imaging 270–271, 277, 521 129 Xe NMR spectroscopy 457–458, 469–470, 476–480 Chemical structure NMR 324–326 elastomers 401–404 NR 433–436 cis-polyisoprene vulcanisate 330–333 Chlorination BR 135–138 chlorocarbenes 136 IR spectra 135–138 Chlorocarbenes, chlorination 136 Chloroprene rubber (CR) crystallinity 91 hydrogenation 134 IR spectra 83, 85, 86, 90, 126, 127, 134 isomers 90–91 Chlorosulfonated polyethylene (CSM) IR spectra 85, 96–97, 99 PA–FTIR spectra 58–60, 102–104 rubber–filler interaction 102–104 self-crosslinking blends 96–98, 99 Cchromatography see high-pressure liquid chromatography Cis–trans isomerisation BR 139, 212–213, 334–335, 336, 338–339 IR spectra 139, 212, 214 polyisoprene 323–324, 328, 338–339, 341, 343, 347, 403, 429–435 see also isomers

614

CK (co-operative kinematics theory) 494– 496 13 C NMR spectroscopy binary blends 340 carbon-black-filled rubbers 341–347 chemical shift 325–326, 328, 524, 543 crosslinking 239, 329–330 crosslink density 329–330, 353, 508– 510 elastomers 402–403 EPM 405–408, 413–414, 508–509 EVA 428–429 1 H–13C heteronuclear coupling 543– 547 high-pressure vulcanisation 339 high-resolution 510 latex state 446–448 linewidths 508–510 magic-angle spinning 329–330, 336, 510, 535–536 NBR 416, 417–418 NR 433–434 PB 414–415, 495–496 peroxide vulcanisates 338–339 polyisoprenes 340, 344, 347, 494–495 naturally occurring 430–432 polyurethane–urea 416, 419 SBR 408–410 silica-filled rubbers 347 silica-filled silicone rubbers 376 sulfur vulcanisates 322 BR 333–338 NR 324–333 SBR 338 swollen state 444–446, 508–510 terminal groups 424–425 thermo-oxidative degradation 348 WISE experiment 543–547 Cohen–Addad–Sotta model 258–259 Cold traps 3, 4–5, 12 Combinatorial chemistry 510

Index Combined rotation and multiple-pulse spectroscopy (CRAMPS) 322 Compatibilisation, PCP/EPDMgDBM 92–94 Composites see carbon-black-filled rubbers Composition see chemical analysis Concentration, IR spectroscopy 82 Conveyor belts, quality control 281 Co-operative kinematics (CK) theory 494–496 Copolymers high-resolution NMR 401–448 chemical composition 402–422 sequence distribution 402, 413–419, 420, 444, 446 IR spectra 92–96 quantitative analysis 88–90 phase-separated systems 547 see also block copolymers Coupling agents 104, 108, 110 Covulcanisation 269–271 CR see chloroprene rubber CRAMPS (combined rotation and multiple-pulse spectroscopy) 322 Crosslinking BR 227, 335–336, 337 characterisation limits 321 13 C NMR 239, 545 crosslink density 234, 271–273 blends 502–510 BR 335–336, 506–507 13 C NMR 329–330, 446, 508–510 double-quantum NMR 541–543 EPDM 224 EPM 508–509 1 H NMR 353, 446, 506–509, 509– 510 double-quantum MAS 541–543 NMR-MOUSE 279–280 EPDM 208–209, 216, 222–225, 237– 239 properties relationships 224–225 filled rubbers 377

inter-crosslink components 505 molecular structure alteration 230 network structure 292–294, 296–307, 345–347 1 H NMR 220, 223–224, 239 NMR imaging models 257–259 reviews 493 NR 226, 326 optical spectroscopy 207–246 peroxide-curing 225–227 co-agents 229–237, 238 polydiene rubbers 210–211, 215–216 polyethylenes 177–178, 179 silane-grafted 174, 175 radiation 156, 177–178 randomly crosslinked chains 302–306, 561 rubbers 209–210 SBR 531–533, 541–542, 545 self-crosslinking blends 96–98, 183, 186 sulfur vulcanisation 323 swelling 491–493 thermoplastic rubbers 168 see also network structure Crystallinity chloroprene rubber 91 EPM 307–308 ethylene–octene copolymer 382–383 polyethylene fibres 459, 460 PVC 183 see also semi-crystalline materials Crystallisation crystallites 354, 584–585 kinetics 308–309 NMR 307–309 PE 169–172 CSDPF (carbon–silica dual phase fillers) 108 CSM see chlorosulfonated polyethylene

615

Spectroscopy of Rubbers and Rubbery Materials Cumyloxy fragments 30, 31 Curing conditions 353, 361, 386 crosslink density 508 curing points 421 (multi) hyphenated TGA 18 network topology 353, 361 physical properties 353, 354 Cyclisation acid-catalysed 420–421 BR 140–142 IR spectra 137, 139–142 Cycloaliphatic epoxide–ENR–GMA, FTIR spectra 100–102 Cyclohexane self-diffusion in BR 497–498 swollen rubbers 266–267, 497–498, 501 Cyclopolymerisation, peroxide-curing coagents 235

D Defects elastomers 264–269 1 H NMR transverse magnetisation relaxation 265–269, 360–366 Deformation low-density polyethylene 170–172 polymers 167–168 rubbery materials 167–168, 557–558, 567–572, 580, 581 sample deformation in NMR imaging 275–276 uniaxially deformed model networks 566–572, 580, 581 Degradation see chemical degradation; thermal degradation Delays alternating with nutations for tailored excitations (DANTE) 509–510 Deproteinisation 433–435

616

Depth profiling 66–70 see also thermal diffusivity Desulfuration 323 Deuterium NMR see 2H NMR Dialkenylsulfides, Fourier transform Raman spectra 217–219, 223 Diallylisophthalate (DAIP) 234–235 1,8-diallyloxyoctane (DAO) 230–232 Diallylphthalate (DAP) 234–235 Diallylterephthalate (DATP) 229, 232–235 Dibutyl maleate (DBM) EPDMgDBM 92–94 grafting 92–94, 172, 174 PEgDBM 92–93 Dicumylperoxide (DCP) 226–227, 232, 233, 338–339 Dicyclopentadiene (DCPD) 207, 228–229 Diene rubbers carbon-black-filled 582, 584 crosslinking 209–210 cyclisation 140 epoxidation 146 hydrosilylation 142, 143 IR spectra 125–126, 127 microstructure 126 oxidation 146 structure 401, 402–403 sulfur vulcanisation 323–324 telechelic 423–424, 439 Differential scanning calorimetry (DSC) and thermogravimetric analysis 1, 2– 3, 5–7, 10–11, 15–17, 29–30 mass spectrometry 2–3, 11, 15–17, 30, 31 Differential thermal analysis (DTA) and thermogravimetric analysis 2–3, 5–6, 10–11, 15–17, 30–33 mass spectrometry 2–3, 6, 11, 15–17, 33–34 Differential thermal gravimetry (DTG) DTG curves 9, 14, 15

Index and thermogravimetric analysis 2–3, 9, 14–16, 36 Diffuse reflectance infrared Fourier transform spectroscopy (DRIFTS) 81, 104, 105 see also Fourier transform infrared spectroscopy 4,4′-diphenylene methane diisocyanate (MDI), polyurethanes 189–190 Dipolar-correlation effect (DCE) 526 Dipolar decoupling (DD) 322, 522–524, 544, 547 Dipole–dipole interactions 13 C NMR 322 decoupling 322, 522–524, 544, 547 double-quantum filtering 540–543 1 H magnetisation-exchange spectroscopy 530 1 H NMR transverse magnetisation relaxation 257–258, 281, 361 multi-quantum filtering 259–264, 538–543 NMR imaging 272, 297, 315, 521– 524 PE 174 soft polymeric matter 306, 310, 315 solid-like properties 257, 520 Disproportionation 424–425 Distortionless enhancement by polarisation transfer (DEPT) 328, 331, 334 Double-quantum NMR 540–543 DSC see differential scanning calorimetry DTA see differential thermal analysis Dynamic mechanical analysis (DMA) 358–9

E Einstein relationship 468 Elasticity 2 H NMR 557–558, 567–568, 575–579

network structure 354, 377–378 NMR 304–305 PDMS 305, 567, 576–579 rubbery materials 557–558, 567–568 soft polymeric matter 291–292, 295, 296, 303 vulcanisation process 321 Elastomer products 265–269 Elastomers ageing 264–265, 274–275, 348 branch-points 426–429, 439 chemical composition 402–422 copolymers 401 defects 264–269 end-groups 404, 405, 422–426, 436– 437, 438 filled 582–584 heterogeneities 264–269 high-resolution NMR 401–456 mechanical load 265 mixing process 264 (multi) hyphenated thermogravimetric analysis techniques 1–47 NMR imaging 247–289 models of transverse relaxation 257– 259 peroxide-curing 225–227 reinforcement 315 rubber–filler interaction 102–108 semi-crystalline 381–383 structure determination 401–404 sulfur vulcanisation 209–210 thermoplastic 168, 188–199, 548– 549, 585–588 vulcanisation 264 129 Xe NMR spectroscopy 457–489 Electromagnetic spectrum 77–78 Electron beam irradiation, IR spectra 156–157, 175–176 Elongation chain structure 296–298, 567–570, 573

617

Spectroscopy of Rubbers and Rubbery Materials PE films 169–170 see also stretching Emulsions, SBR 62–63 End-groups, NMR 404, 405, 422–429, 436–437, 438 ENR see epoxidised natural rubber Entropy reduction 296–297, 345 EPDM see ethylene–propylene–diene terpolymer EPM see ethylene–propylene copolymers Epoxidation, IR spectra 146 Epoxidised natural rubber (ENR) hydrogenation 132, 134 IR spectra 85, 100–102, 104–108, 134 see also natural rubber Epoxidised soyabean oil (ESO) 187 Epoxy titration method 405–406 Ethylene–methacrylic acid copolymer, zinc salt 151–153, 154–155 Ethylene–octene copolymer, crystallinity 382–383 Ethylene–propylene copolymers (EPM) 13 C NMR 405–408, 413–414, 508– 509 crosslink density 508–509 crystallinity 307–308 IR spectra 89, 90, 180, 181 NMR 325, 405–410, 413–414 PA–FTIR spectra 64–65 peroxide-curing 230–231, 232–234, 235, 236 polypropylene blends 465–474, 482 properties 207 propylene determination 90, 405–410, 414, 469 sequence distribution 413–414, 424– 425 129 Xe NMR 465–474 Ethylene–propylene–diene terpolymer (EPDM) carbon-black-filled 369–370, 372,

618

373, 463, 465, 474–480 chain length distribution function 365 crosslinking 208–209, 216, 222–225, 237–239 1 H NMR transverse magnetisation relaxation 386–387 IR spectra 84, 89, 93 grafting 92–94, 95, 156–157 ionomeric blends 151–153, 154–155 termonomer content 90, 91 mechanical properties 224–225, 238 network structure 223–225 NMR imaging 254–255 oil-extended 366–367 peroxide-curing 209, 227–237 co-agents 230, 238 third monomer effects 227–229, 238 polypropylene blends 482–485 production 207–208 properties 208 SEC–NMR 410–413 sulfur vulcanisation 208–209, 216– 225, 238, 365, 386–387, 422 5-ethylidene-2-norbornene conversion 219–221, 224, 238 thermal decomposition 23–24, 30, 31 129 Xe NMR 458, 463–464, 474–480, 482–486 Ethylene vinyl acetate (EVA) branch-points 428–429 electron beam irradiation 156–157 IR spectra 172, 173 NMR 416, 419, 420, 428–429 polyethylene copolymer 172, 173 2-ethylhexyl acrylate (EHA), 1H NMR transverse magnetisation relaxation 358– 359, 361–363 5-ethylidene-2-norbornene (ENB) 207, 216–217, 219–221, 224, 238, 410–413 EVA see ethylene vinyl acetate Evolved gas analysis, (multi) hyphenated TGA 15, 17, 18

Index Evolved gas profiles (EGP) 9, 21 Extinction coefficient 82

F Far infrared region 78 Fibres, PE 459, 460 Filled rubbers chain grafting 379–381 elastomers 582–584 filler clusters 265–266 1 H NMR transverse magnetisation relaxation 368–379 2 H NMR 582–584 network structure 377–378 rubber–filler interaction 102–108, 368, 371, 372, 374, 376, 380, 474– 475 see also carbon-black-filled rubbers; silica-filled rubbers; silica-filled silicon rubbers Films Kapton 70 PE 68–69, 169–170, 174–178 Fire resistant compounds 92–93 Food industry, radiation processing 174 Fourier transform infrared spectroscopy (FTIR) 81, 211 carboxylate ionomers 147–149 peroxide-curing co-agents 232, 233, 234–236, 237, 238 photoacoustic 49–76 polyethylenes aluminium laminates 178–179 low-density 170–172, 176–178 medium-density 178, 179 PVC blends 181–182 sulfonated ionomers 149–151 sulfur-vulcanised NR 214 and thermogravimetric analysis 2–4, 6–10, 11–12, 15–18, 19–25, 34–35

see also diffuse reflectance infrared Fourier transform spectroscopy ; infrared spectroscopy Fourier transform NMR 248–251, 508, 528–529, 544 Fourier transform Raman spectroscopy 211 dialkenylsulfides 217–219 sulfur vulcanisation 217–223, 224 see also Raman spectroscopy free radicals carbon-black-filled rubbers 368, 369 macroradical side reaction suppression 230–232 sulfur vulcanisation 213, 335 Frequencies 77–78 Larmour frequencies 248–249, 295, 297, 462, 521, 523, 526 stretching 79–80 vibrational 78–80 FTIR see Fourier transform infrared spectroscopy; photoacoustic Fourier transform infrared spectroscopy Functional groups analysis 80 formation during milling 109, 111

G γ-irradiation low-density polyethylene film 176–178 NR 339 PIB 426 polyisoprene 503–504 polymers 153 vulcanisation 339 Gas chromatography (GC) and thermogravimetric analysis 3–5, 11–12, 15–17 mass spectrometry 3–4, 11, 15–17, 34–35 Gaskets, NMR imaging 265–266

619

Spectroscopy of Rubbers and Rubbery Materials Gas phase, 129Xe NMR spectroscopy 458– 459 Gated high-power decoupling (GHPD) 334 GC see gas chromatography Gels 291, 292, 299, 301–305, 310, 447 gelation 302–304, 307 stretched 306 swelling 491–492 Glass-to-rubber transition temperature (Tg) 168, 279–280, 493, 520, 524–525, 557, 562 Glassy materials, log (shear) modulus 167 Glycidyl methacrylate (GMA) 100–102 Goodyear 209 Gotlib–Fedotov–Schneider model 258, 259 Gradient echoes NMR imaging 253–255, 265, 271, 462–463, 496–497, 498–500 129 Xe NMR spectroscopy 462–463 Grafting dibutyl maleate 92–94, 172, 174 EPDM 92–94, 95, 156–157 filler surface 379–381 IR spectra 92–96, 156–157, 172, 174 NMR 421 polyethylene 92–94, 172, 174, 175– 176, 178 PVC 188 radiation induced 178 vinyl silanes 174 Guayule rubber, isomers 88 Gutta percha, isomers 88, 403, 429, 431

H Hahn-echo imaging 253–257, 258–259, 262, 265, 266–267, 268, 269–271, 274– 275, 277–280, 282, 503, 527 Halogenation, IR spectra 135–138

620

Hamiltonians 521, 522, 540 Helium atmosphere, photoacoustic Fourier transform infrared spectroscopy 53–54 Herschel, Sir William 77 Heterogeneity elastomers 264–269 1 H NMR transverse magnetisation relaxation 265–269, 360–366 molecular-scale 360 morphological 361 network structure 353, 355, 360–366 spatial 273, 353, 361 Heteronuclear correlation experiments 543–547 Hevea rubber cyclisation 142 isomers 88, 403 NMR 436 High-density polyethylene (HDPE), IR spectra 194–198 High-performance liquid chromatography 404 High-pressure liquid chromatography 331–332 High-pressure vulcanisation (HPV) 339 High-resolution NMR 13 C MAS NMR 510 elastomer structural characterisation 401–456 applications 436–448 multinuclear 436–443 vulcanisate chemical characterisation 321–352 1 H magnetisation-exchange spectroscopy 529, 530–535 1 H NMR relaxation butyl rubber 426–427 CRAMPS techniques 322 crosslinking 220, 223–224, 239, 299– 306

Index crosslink density 353, 506–508, 509– 510 double-quantum spectroscopy 540– 543 elastomers 402–403 EPDM 410–413 1 H–13C heteronuclear coupling 543– 547 linewidths 506–508, 509–510 multi-quantum spectroscopy 259–264, 538–543 nanometer scale probe 291 NOESY–MAS spectroscopy 535–538 NR 433, 435 PB 314 PDMS rubber bands 262–263 pseudo-solid spin-echoes 298–299, 310 water/ice 321–322 WISE experiment 543–547 1 H NMR transverse magnetisation relaxation blends 384 chain dynamics 313, 355–360, 361, 371, 374–376, 379 crosslink density 257–259, 502–506 decay curves 363–364, 366, 382, 502– 505 emulsions 385 EPDM sulfur vulcanisation 386–387 filled rubbers 368–379 carbon-black-filled rubbers 265–266, 369–374, 475, 504 silica-filled rubbers 378–381 silica-filled silicone rubbers 374–378 multi-quantum spectroscopy 260–262 network structure analysis 355–379 chain entanglements 309–315, 366– 367, 377–378, 505–506 heterogeneity and defects 265–269, 360–366 molecular mass distribution 364–366

network density 356–360, 367 real-time experiments 386–387 NR 267–268, 279 quality control 281, 387–388 real-time experiments 385–387 residual dipolar coupling 524–526, 527 rubbery material characterisation 353–400 SBR 267–268, 279–280 semi-crystalline elastomers 381–383 solvent effects 361–363, 499 swollen state 354, 361–363, 378, 500–506 temperature dependence 276, 356– 357, 375, 376 theory 248–249, 252, 255–259, 295, 297, 299 thermo-oxidative ageing 274–275 viscoelastic materials ionic 383–384 molecular dynamics 526–527 viscoelastic properties 365–366 vulcanisation process 273–274 see also NMR imaging; NMRMOUSE 2 H NMR chain extension 558 deuterated probes 579–581 elastic property correlation 575–579 elastomers 436–437, 438 filled 582–584 thermoplastic 548–549, 585–588 gels 306 line shape analysis 527, 561–563, 566–575, 580–581 mean-field effect 569–571, 575 model systems 561–579 molecular dynamics 559–561 orientation effects 560–561, 569, 571, 573–575, 579–580

621

Spectroscopy of Rubbers and Rubbery Materials PB 436–437, 571–572, 577, 580–581, 582, 584, 587–588 PDMS 263, 527, 528, 561–572, 574– 575, 576–579, 584, 585 pseudo-solid spin-echoes 561–563 quadrupolar interactions 527, 528, 559, 564–565, 567–571, 588 rubbery materials 557–596 semi-crystalline polymers 584–585 silica-filled silicone rubbers 584 swollen state 580–581 theory 559–561 time scale of motions 563–566 uniaxially deformed model networks 566–572, 580, 581 viscoelastic materials 522, 527 Homogeneous polarisation field (B0) 247, 250, 257–258, 277, 282, 368 Hydroboration, IR spectra 137, 144 Hydrocarboxylation, IR spectra 146 Hydroformylation IR spectra 137, 144–145 polypentenamer 144–145 Hydrogen abstraction 424–425 Hydrogenation BR 129–131 CR 134 degree of hydrogenation 128–129 ENR 132, 134 IR spectra 126, 127–134, 137, 144– 145 NBR 126, 127–129, 421 NR 131–132 SBR 131–132 Hydrogen bonding PB 587–588 plasticised PVC 168 polyureas 192, 193 polyurethanes 100, 190–192 PVC 174 rubbery materials 354

622

silica-filled rubbers 378 thermoplastic elastomers 548 Hydroquinone bis (β-hydroxyethyl) ether (HQEE) 189–192 Hydrosilylation BR 142–144 IR spectra 137, 142–144 Hydroxylation, IR spectra 137 2-hydroxyl ethyl methacrylate (HEMA) 188 Hypalon, self-crosslinking blends 97 Hyphenated techniques see (multi) Hyphenated thermogravimetric analysis (TGA) techniques

I IIR see butyl rubber Impurities (multi) hyphenated TGA 27 NMR small signals 422, 477 Infrared spectroscopy (IR) 210–211 adhesion 111–113 applications 82–113, 125–165 blends 92–98, 111–113, 180, 182–183 bromination 137 carbon black 198–199 chemically modified rubbers 125–165 chlorination 135–138 cyclisation 137, 139–142 degradation 113 halogenation 135–138 hydroboration 137, 144 hydrocarboxylation 146 hydroformylation 137, 144–145 hydrogenation 126, 127–134, 137, 144–145 hydrosilylation 137, 142–144 hydroxylation 137 ionomers blends 151–153, 154–155

Index formation 138, 147–151 sulfonated 149–151, 198–199 irradiation 153, 156–157, 174–176, 177 isomerisation 139, 212, 214 milling 108–109, 111 oxidation 137, 146 phosphonylation 137, 146–147 polyamide-12 192–193 polyethylenes 93, 168–180, 181 low-density 83, 169, 170–172, 176– 178 medium-density 178, 179 polyureas 192, 193 polyurethanes 91–92, 98, 100–102, 188–194 PVC 180–188 quantitative analysis 81–82, 88–90 reverse engineering 113 rubbers 77–124 ASTM D3677 83 peak assignments 83–86 rubber–filler interaction 102–108 rubbery materials 167–205 characterisation 125–165 sample preparation 80–81, 83 silica 198–199 spectrometers 77 sulfonation 138, 149–151 sulfur-vulcanised NR 214 theory 77–80 thermoplastic elastomers 188–199 weathering 138, 153 zinc stearate 198–199 see also Fourier transform infrared spectroscopy (FTIR) Interface properties 2 H NMR 584 NMR imaging 269–271, 273 rubbery materials 354, 559 rubber–filler 102–108, 368, 371, 372,

374, 376, 380, 474–475 Interfaces, coupled instruments 6, 10 Interlaboratory reproducibility 13–14 International Confederation for Thermal Analysis (ICTA) recommendations 13 International Institute of Synthetic Rubber Producers (IIRP) 64 Interpenetrating networks (IPN) 92 Ionic bonding rubbery materials 354 viscoelastic materials 383–384 Ionomer formation blends 151–153, 154–155 1 H NMR transverse magnetisation relaxation 383–384 IR spectra 138, 147–151 polyurethanes (PU) 194–195 sulfonated 149–151, 198–199 IPN (interpenetrating networks) 92 iPP (isotactic polypropylene) 465–474, 482–485, 585 IR see infrared spectroscopy Irradiation see electron beam irradiation; γ-irradiation; radiation Isomers balata 88 BR 83, 88, 90, 139, 212–213, 403 CR 90–91 cyclisation 140–142 NR 88, 403 PB 414–415 polyisoprene 88, 323–324, 328, 338– 339, 341, 343, 347, 403 see also cis–trans isomerisation; cispolyisoprene Isotactic polypropylene (iPP) 465–474, 482–485, 585

J J-resolved spectroscopy 543

623


K

M

Kapton film, PA–FTIR spectra 70 Kinetics crystallisation 308–309 gelation 307 (multi) hyphenated TGA 18 k-space coordinates 252–253, 282 Kuhn segments 259

Macroradical side reaction suppression 230–232 Magic-angle spinning (MAS) 522–524 13 C NMR 322, 329–330, 336, 510, 535–536 1 H double-quantum spectroscopy 540–543 NOESY–MAS spectroscopy 535–538 WISE experiment 544, 545–547 Maleation 92, 421 MAS see magic-angle spinning Mass spectrometry quadrupole 10 and thermogravimetric analysis 2–4, 7, 8, 10–12, 15–17, 18–19, 25–29, 35–36 differential scanning calorimetry 2–3, 11, 15–17, 30, 31 differential thermal analysis 2–3, 6, 11, 15–17, 33–34 gas chromatography 34–35 time-of-flight 10, 28, 36 MDI (4,4′-diphenylene methane diisocyanate) 189–190 Mechanical load, NMR imaging 265 Mechanical properties EPDM 224–225, 238 filled rubbers 368 network structure 353, 354, 366–367 heterogeneity 265–269, 360–366 oil-extended rubbers 366–367 peroxide-curing co-agents 230, 236– 237 Medium-density polyethylene (MDPE), IR spectra 178, 179 5-methylidene-2-norbonene (MNB), EPDM peroxide-curing effects 228 Methyl methacrylate (MMA), polyethylene graft copolymers 175–176 Microstructure

L Laminates, PA–FTIR spectra 70 Larmour frequencies 248–249, 295, 297, 462, 521, 523, 526 Latex state high-resolution NMR 446–448 NR 433 PB 279–280, 447 suspensions 306 LDPE see low-density polyethylene Least squares linear regression analysis 60–65 7 Li NMR 437, 439–440 Line shape analysis, 2H NMR 527, 561– 563, 566–575, 580–581 Liquid-like properties 520, 527, 533, 562–563, 564 see also solution NMR Liquid–lattice theory 491–492 Lithium initiators 436, 437, 439 Living polymers 436–439 Loaded polymers 293–294, 315–317 Long-chain molecules 313–315 Low-density polyethylene (LDPE) ethylene vinyl acetate copolymer 172, 173 IR spectra 169, 170–172, 188, 189 γ-irradiation 176–178 NR blends 188, 189 sample preparation 83

624

Index 13

C NMR 340 diene rubbers 126 PB 403 polyisoprene 403 Milling 108–109, 111 Mineral aggregates, chain adsorption 293–294, 315–317 Mirror velocity 51, 66–68 Mixing, elastomers 264 Mobile universal surface explorer (NMRMOUSE) 268–269, 277–284, 388 Molecular dynamics 524–528, 530–538 2 H NMR 559–561 slow motions 526, 529, 548–549, 564–565 see also chain dynamics Molecular properties, rubbery materials 557–558 Molecular structure conformation 504 diene rubbers 126 PB 83, 87 peroxide-cured elastomers 230 polyisoprene 87 see also chain structure Molecular vibration 78–80 Monte Carlo method 573 Mooney–Rivlin coefficients 578–579 MOUSE see NMR-MOUSE Multidimensional NMR techniques principle 528–529 viscoelastic material characterisation 519–555 (Multi) hyphenated thermogravimetric analysis (TGA) techniques applications 14–35 calibration 13–14 elastomer characterisation 1–47 future prospects 35–36 quality control 14, 28, 29 sample requirements 13

software packages 6 Multinuclear high-resolution NMR 436– 443 Multi-quantum NMR 259–264, 538–543 Mushroom rubber 431, 432

N Nanometer scale probe 291–320 Natural rubber (NR) biosynthesis 402, 403 13 C NMR 324–333 crosslinking 226, 326, 329–330 crosslink density 506–508 deproteinisation 433–435 γ-irradiated 339 1 H NMR 433, 435 1 H NMR transverse magnetisation relaxation 267–268, 279 IR spectra 83, 84, 86, 131–132 epoxidation 146 hydrogenation 131–132 isomers 88, 403, 429 latex state 433 low-density polyethylene blends 188, 189 NMR-MOUSE 279–280 PA–FTIR spectra 56–57 peroxide-curing 226, 338–339 small molecule diffusion 499–500 structure 88, 433–435 sulfur vulcanisation 214, 323–333, 334 accelerated 327–333 carbon-black-filled 341–344, 582, 583 covulcanisation 269–271 unaccelerated 327 swollen 506–508 terminal groups 436 see also epoxidised natural rubber; polyisoprene

625

Spectroscopy of Rubbers and Rubbery Materials NBR see nitrile–butadiene rubber Network structure chain structure 293–294, 296–307, 333, 355 chemical analysis 354–355 13 C NMR 339, 433–434 defects 360–366 disentanglement 361 end-linked 299, 561, 563, 575, 577–579 entanglements 309–315, 354, 360, 366–367, 377–378, 505–506, 561 EPDM 223–225 filled rubbers 377–378 heterogeneity 264–269, 353, 354, 360–366 1 H NMR transverse magnetisation relaxation 355–379, 386 molten high polymers 292, 309–314, 502–504 ionic viscoelastic materials 383–384 loaded polymers 293–294 mechanical property effects 353, 354, 360–366 molecular mass distribution 364–366 molten high polymers 292, 309–314, 502–504 network distribution function 294– 295, 299–301 NMR imaging 271–273 models 257–259 pseudo-solid spin-echoes 298–299, 310, 561–563 oil-extended rubbers 366–367 PDMS 566–572 properties relationships 224–225, 236–237 randomly crosslinked 302–306, 561 sulfur vulcanisation 323 swollen rubbers and blends 502–510 temporary networks 310–312 thermo-oxidative degradation 347–348

626

topology 353, 361, 378 uniaxially deformed model networks 566–572 see also crosslinking nitrile–butadiene rubber (NBR) acrylonitrile determination 61–62, 88– 90 carbon-black-filled 504 13 C NMR 416, 417–418 crosslink density 506 hydrogenation 126, 127–129, 421 IR spectra 83, 85, 88, 89, 111–112, 126–129 PA–FTIR spectra 61–62 PVC blends 111–113, 183–186 see also carboxylated nitrile rubber NMR angular dependence 521, 522, 525, 560–561 anisotropic effects 322, 519–522, 524, 529, 558, 559–560 branch-points 426–429, 439 chemical composition 402–422 accuracy 404–410 distribution 410–413 chemically modified structures 419– 422 chemical shift 270–271, 277, 325– 326, 328, 457–458, 469–470, 476– 480 chemical structure 324–326, 330–333, 401–404 correlation 336, 525–526, 529, 538 dipolar-correlation effect 526 heteronuclear correlation experiments 543–547 cross-polarisation 322, 470, 474 crystallisation 307–309 dipolar decoupling 322, 522–524, 544, 547 end-groups 422–426, 436–437, 438 EPM 307–308

Index field cycling technique 526 high-resolution 401–456 latex state 306, 446–448 7 Li NMR 437, 439–440 linewidths 506–510 linkage effect 302 low-resolution 248, 277 magic-angle spinning 322, 329–330, 336, 510, 522–524 multidimensional techniques 519–555 nanometer scale probe 291–320 orientation dependence 281, 521–522, 525, 527, 560–561, 569, 571, 573– 575, 579–580 31 P NMR 442 pseudo-solid spin-echoes 298–299, 310, 527, 548, 561–563 quadrupolar interactions 522, 527, 528, 559, 564–565, 567–571, 588 quantitative analysis 405–410, 541 residual dipolar coupling 524–526, 527, 530–531, 533–535, 538–547 sequence distribution 402, 413–419, 420, 430, 444, 446 29 Si NMR 440–441 size exclusion chromatography 404, 410–413 small signal assignment 422 29 Sn NMR 442–443, 444 soft polymeric matter 291–320 solid-state methods 321–322 solution methods 321–322 theory 295 two-dimensional 519–520 see also 13C NMR spectroscopy; dipole–dipole interactions; highresolution NMR; 1H NMR relaxation; 1 H NMR transverse magnetisation relaxation; 2H NMR; multidimensional NMR techniques; 129 Xe NMR spectroscopy

NMR imaging applications 264–276 blends 264, 271, 272, 384 BR 500–501 calibration 256–257 chemical shift 270–271, 277, 521 contrast 250, 253–264, 270–271, 558 multi-quantum filtering 259–264 relaxation-time contrast 255–257, 271–276 susceptibility contrast 254–255 coordinates 252–253 CPMG echoes 278–280 crosslink density 271–273, 279–280, 446 defects 264–269 elastomers 247–289 Fourier NMR 248–251, 508, 528– 529, 544 Hahn echoes 253–257, 258–259, 262, 265, 266–267, 268, 269–271, 274– 275, 277–280, 282, 503, 527 heterogeneities 264–269 homogeneous polarisation field (B0) 247, 250, 257–258, 277, 282 interface detection 269–271, 273, 559, 584 Larmour frequencies 248–249, 295, 297 low-resolution 248, 277, 384 magnetic field gradients 250–254, 273, 280, 282, 369 gradient echo imaging 253–255, 265, 271, 462–463, 496–497, 498–500 magnetisation 248–249, 252, 255– 256, 258, 295, 297, 299 models 257–259 Cohen–Addad–Sotta model 258–259 Gotlib–Fedotov–Schneider model 258, 259 NR 499–500 polyisoprene 500

627

Spectroscopy of Rubbers and Rubbery Materials precession phase 248–250 pulsed 250, 253–254, 522, 523, 528, 530, 543–544 sample deformation 275–276 single-point imaging 267–268, 282, 283 space encoding 253, 255–256, 282 spatial resolution 248, 251–253, 255, 276–277 SPRITE technique 267–268 swollen rubbers 499–502 theory 248–264 vulcanisation process 264, 273–274, 303–304 NMR-MOUSE (mobile universal surface explorer) 268–269, 277–284, 388 NMR relaxation techniques elastomers 247–248, 271–276 relaxation-time contrast 255–257 soft polymeric matter 295 see also 1H NMR relaxation; 1H NMR transverse magnetisation relaxation NOE (nuclear Overhauser effect) 414, 535–538 Non-Newtonian liquids 292 NR see natural rubber Nuclear magnetic resonance see NMR Nuclear Overhauser effect (NOE) 414, 535–538 Nylon 6, acrylate rubber blends 194

O Oil-extended rubbers, network structure 366–367 Opacity 51 Optical absorption coefficient 51 Optical spectroscopy crosslinking 207–246 sulfur vulcanisation 211, 214 see also infrared spectroscopy; NMR; Raman spectroscopy

628

Orientation effects 2 H NMR 560–561, 569, 571, 573– 575, 579–580 NMR-MOUSE 281 PA–FTIR spectra 71 polyethylene films 169–170 soft polymeric matter 291 viscoelastic materials 527 Outgassing, plasticisers 23–24 Oxidation IR spectra 137, 146 photo-oxidation 68, 70 polyethylene film 68–69 see also thermo-oxidative degradation Ozone treated polyethylene film, PA– FTIR spectrum 68–69

P PA–FTIR see photoacoustic Fourier transform infrared spectroscopy PB see polybutadiene PDMS see poly(dimethylsiloxane) PE see polyethylene Percolation 302–304, 315 Peroxide-curing 210, 225–226, 338–339 BR 227, 338–339 co-agents 229–237, 238 EPDM 209, 227–237 third monomer effects 227–229, 238 EPM 230–231, 232–234, 235, 236 NR 226, 338–339 PB 227 polydiene elastomers 225–227 polyisoprene 339, 506 scorch retarders 237 Phase-separated systems 547 Phenol–formaldehyde resins (resols) 209 Phosphonylation, IR spectra 137, 146–147 Photoacoustic detectors 50–51, 53 Photoacoustic effect 49–52

Photoacoustic Fourier transform infrared spectroscopy (PA–FTIR) 81 background spectrum 53 baseline constructions 64–65 depth profiling 66–70 helium atmosphere 53–54 history 49–50 instrumentation 49–51, 52–55, 57 polarised 71 powders 54–55 quantitative analysis 60–65 rubbers and related materials 49–76 samples 54–55 sensitivity 53–54 software 55, 57 surface analysis 66–70 theory 50–52 Photoageing polyurethanes 194 PVC 187–188 see also ageing Photons 77–78 Photo-oxidation 68, 70 Physical networks see network structure Physical properties, cured materials 354 Planck’s constant 77 Plasticisers outgassing 23–24 PVC hydrogen bonding 168 Plasticity 167–168 13 P NMR 442 Polarised photoacoustic Fourier transform infrared spectroscopy 71 Polyamide 6,66,610 copolymer (LPA), IR spectra 94–95 Polyamide-12 (PA-12), IR spectra 192– 193 Polyamide–polyether (PA–PEG) block copolymers 585–587 Polybutadiene (PB) acid-catalysed cyclisation 420–421

carbon-black-filled 582, 584 C NMR 414–415, 495–496 end-groups 436–437 FTIR spectra 88 high-pressure vulcanisation 339 1 H NMR 314 2 H NMR 436–437, 571–572, 577, 580, 581, 582, 584, 587–588 IR spectra 83–84, 87–88 isomers 83, 88, 90, 139, 403, 414– 415 latex state 279–280, 447 peroxide-curing 227 29 Sn NMR 442–443, 444 structure 83, 87 telechelic 423–424 temporary network 311 Polybutylacrylate 437, 439 Poly-σ-caprolactone (PCL), PVC blends 182–183 Polychloroprene (PCP) IR spectra 92–94 NMR 416, 444, 445 Polydiene rubbers crosslinking 210–211, 215–216 peroxide-curing 225–227 sulfur vulcanisation 211–216 mechanism 214–216 Poly(dimethylsiloxane) (PDMS) adsorption on mineral aggregates 315–316 crosslink density 505, 561 elasticity 305, 567, 576–577 2 H NMR 527, 528, 561–572, 574– 575, 576–579, 584, 585 NMR imaging 256–257, 262–263, 305 silica-filled silicone rubbers 374–381 uniaxially deformed network 566–572 Polyester–polyurethane, chemical degradation 442 13

629

Spectroscopy of Rubbers and Rubbery Materials Polyethylene (PE) aluminium laminates 178–179 fibres 459, 460 films 68–69, 169–170, 174–178 2 H NMR 585 IR spectra 93, 168–180, 181 chain branching 168–169 crosslinking 174, 175, 177–178, 179 crystallisation 169–170 elongation 169–170 grafting 92–94, 172, 174, 175–176, 178 irradiation 176–178 orientation effects 169–170 peroxide-cured 178, 179 pipe welding 178, 179 single crystals 180 surface oxidation 68–69 129 Xe NMR spectroscopy 459, 460 see also high-density polyethylene, low-density polyethylene, mediumdensity polyethylene Poly(ethylene glycol) diacrylate (PEGDA), 1 H NMR transverse magnetisation relaxation 358–359, 361–363 Polyisobutylene (PIB) γ-irradiation-induced scission 426 1 H NMR 426–427 Polyisoprene acid-catalysed cyclisation 420–421 cis–trans isomerisation 323–324, 328, 338–339, 341, 343, 347, 403, 429– 435 13 C NMR spin–lattice relaxation 494– 495 crosslink density 506, 543 end-groups 436–437, 438 γ-irradiation 503–504 1 H NMR double-quantum spectroscopy 543 transverse magnetisation relaxation 503–504

630

2

H NMR 436–437, 438 IR spectra cyclisation 142 epoxidation 146 isomers 88, 131 7 Li NMR 439–440 multi-quantum NMR 260–261 naturally occurring 429–436 NMR 416, 500 peroxide-curing 339 structure 87, 88 sulfur vulcanisation 324–325 swollen 500 telechelic 423–424 129 Xe NMR 485–486 see also natural rubber cis-polyisoprene (IIR), sulfur vulcanisate 330–333, 340, 343–344, 347 Polyketones 146 Polymer blends see blends Polymeric networks see network structure Polymers aggregated 294 amorphous 460, 585 chemical modification 125 crosslinking 493 defects 364 deformation 167–168 γ-irradiation 153 gels 491–492 living 436–439 loaded 293–294, 315–317 molten 292, 309–314, 502–504, 562, 564 orientation effects 71, 169–170 quantitative analysis 60–65 semi-crystalline 292–293, 307–309, 381–382, 481, 584–585 solutions 494–496 thermogravimetric analysis samples 12–13

viscoelastic deformation 167–168 129 Xe NMR spectroscopy 459–461, 481–482 see also copolymers; soft polymeric matter Poly(methyl methacrylate) (PMMA) quantitative NMR 404–405 129 Xe NMR spectroscopy 460–461 Polypentenamer (PPA), hydroformylation 144–145 Polyprenols 429–433 Poly-β-propiolactone (PPL), PVC blends 182–183 Polypropylene (PP) blends and block copolymers 465– 474, 482–485 isotactic 465–474, 482–485, 585 phosphonylation 137, 146–147 Poly(propylene oxide) (PPO) networks, 1 H NMR transverse magnetisation relaxation 356–357 Polystyrene, PA–FTIR spectra 54, 55 Polyureas hydrogen bonding 192, 193 IR spectra 192, 193 Polyurethanes (PU) degradation 194 2 H NMR 587 hydrogen bonding 100, 190–192 ionomer formation 194–195 IR spectra 91–92, 98, 100–102, 188– 194 phase-separated systems 547 31 P NMR 442 Polyurethane–urea 416, 419 13 C NMR 416, 419 Polyvinyl chloride (PVC) blends 180, 182–187 NBR 111–113, 183–186 dechlorination 138, 186 grafting 188

hydrogen bonding 174 IR spectra 180–188 photoageing 187–188 plasticised 168 substitution reactions 188 thermo-oxidative degradation 185– 187 Powders, photoacoustic Fourier transform infrared spectroscopy 54–55 PP see polypropylene Principle component analysis (PCA) 36, 408, 410 Processing elastomers 264–276 EPDM 207–208 quality control 281 Product development, (multi) hyphenated TGA 18 Propylene, determination in EPM 90, 405–410, 414 Propylene–ethylene copolymers see ethylene–propylene copolymers Proton NMR magnetisation relaxation see 1H NMR relaxation; 1H NMR transverse magnetisation relaxation Pseudo-solid spin-echoes 298–299, 310, 527, 548, 561–563 PTA (pulse thermal analysis) 19 PU see polyurethanes Pulsed field gradient NMR 496–497, 498–499 129 Xe NMR 462–463, 480–482, 483– 484, 485 Pulse thermal analysis (PTA) 19 PVC see polyvinyl chloride Pyrolysis methods 35

Q Quadrupolar interactions 522, 527, 528, 559, 564–565, 567–571, 588

631

Spectroscopy of Rubbers and Rubbery Materials Quadrupole mass spectrometry (QMS) 10 Quality control 1 H NMR transverse magnetisation relaxation 281, 387–388 (multi) hyphenated thermogravimetric analysis techniques 14, 28, 29 NMR-MOUSE 281 Quantitative analysis automated analysis 408 copolymers 88–90 IR spectroscopy 81–82, 88–90 NMR 405–410, 541 PA–FTIR 60–65 polymers 60–65 TGA 18–19

R Radiation IR spectra 153, 156–157, 174–176, 177 vulcanisation 339 see also electron beam irradiation; γirradiation Raman spectroscopy 210–211 sulfur vulcanisation 211–213, 214, 238 see also Fourier transform Raman spectroscopy Reactivity, (multi) hyphenated TGA 18 Real-time 1H NMR transverse magnetisation relaxation 385–387 Recycling, thermoplastic elastomers 188 residual dipolar coupling 524–526, 527, 530–531, 533–535, 538–547 Resols (phenol–formaldehyde resins) 209 Reverse engineering 113 Rouse model 313 Rubber blends see blends Rubbers crosslinking 209–210, 493

632

high-resolution 13C MAS NMR 510 IR spectroscopy 77–124 photoacoustic Fourier transform infrared spectroscopy 49–76 quantitative analysis 81–82, 88–90 self-diffusion 496–499 swelling 491–493, 502–510 vulcanisates 209–210 see also natural rubber Rubbery materials chemically modified 125–165 customised properties 557 deformation 167–168, 557–558, 567– 572 1 H NMR transverse magnetisation relaxation 353–400 2 H NMR 557–596 IR spectroscopy 167–205 characterisation 125–165 molecular properties 557–558 swelling 491–518, 580 Rubber–filler interactions 102–108, 368, 371, 372, 374, 376, 380, 474–475, 584

S Samples IR spectroscopy 80–81, 83 NMR imaging 275–276 photoacoustic Fourier transform infrared spectroscopy 54–55 thermogravimetric analysis 12–13 SBR see styrene–butadiene rubbers Scorch retarders, peroxide-curing 237 Seals see gaskets Self-crosslinking blends, IR spectra 96– 98, 183, 186 Self-diffusion rubbers 498–499 small molecules 496–500 129 Xe NMR 463, 464, 480–487

diffusion time 485–487 Semi-crystalline materials elastomers 381–383 polymers 292–293, 307–309, 381– 382, 481, 584–585 see also crystallinity Sequence distribution 402, 413–419, 420, 424–425, 430, 444, 446, 495–496 Shear modulus 167–168, 530 Short-chain molecules 312 Silica IR spectra 198–199 PA–FTIR spectra 56–57 Silica-filled rubbers carbon–silica dual phase fillers 108 chain adsorption 315–316, 375–377, 378–379 grafting 379–381 1 H NMR transverse magnetisation relaxation 378–381 PDMS 315–316 cis-polyisoprene 347 rubber–filler interaction 106, 108, 109, 110, 584 SBR 66–67 sulfur vulcanisation 347 Silica-filled silicon rubbers 1 H NMR transverse magnetisation relaxation 374–378 2 H NMR 584 Silicone methyl rubber, IR spectra 86 Siloxanes, 29Si NMR 440–441 Single-point ramped imaging with T1 enhancement technique (SPRITE) 267–268 29 Si NMR 440–441 Size exclusion chromatography (SEC) 404, 410–413 29 Sn NMR 442–443, 444 Soft polymeric matter NMR 291–320 NMR imaging 247, 282

Software (multi) hyphenated thermogravimetric analysis 6 photoacoustic Fourier transform infrared spectroscopy 55, 57 Solid-like properties 257, 520, 527, 533 Solid state NMR see 13C NMR Spectroscopy; high-resolution NMR; 1H NMR relaxation Solution NMR 321–322, 519, 522 13 C NMR 494–496, 508–509 2 H NMR 559 multi-quantum NMR 538 NOESY–MAS spectroscopy 535 swollen state 443–446, 500 see also liquid-like properties Solvent effects 13 C NMR 444, 445 1 H NMR transverse magnetisation relaxation 361–363, 499 2 H NMR 580–581 see also swollen state Spatial heterogeneity crosslink density 273 network structure 353, 361, 558 Spatially resolved NMR 276–277, 558 Spin–lattice relaxation 310, 494–496 Spin–spin interactions 295, 297, 298– 299, 300–301, 314, 520–524 manipulation 522–524 spin diffusion 547 SPRITE (single-point ramped imaging with T1 enhancement) technique 267–268 Stability see thermal stability Standardisation interlaboratory reproducibility 13–14 see also ASTM standards Storage modulus 358–359 Stress/strain effects 2 H NMR 573–575, 579–580, 589 NMR imaging 265

633

Spectroscopy of Rubbers and Rubbery Materials strain distribution 263–264 Stretching gels 306 polyethylene fibres 459, 460 rubbery materials 567–570, 574 see also elasticity; elongation Stretching vibrations 79–80 Structure see chain structure; chemical structure; microstructure; molecular structure; network structure Styrene/ethylene–butylene/styrene (SEBS), sulfonation 149–151 Styrene–butadiene rubbers (SBR) carbon-black-filled 24–25, 372, 374, 582 13 C NMR 336 covulcanisation 269–271 crosslinking 531–533, 541–542, 545 emulsion SBR 62–63 1 H NMR double-quantum MAS spectroscopy 540–542 magnetisation-exchange spectroscopy 530–533 NOESY–MAS spectroscopy 536–538 1 H NMR transverse magnetisation relaxation 267–268, 279–280 2 H NMR 582 hydrogenation 131, 132 IR spectra 83, 84, 88–89, 127, 131, 132 NMR 408–410, 522, 523 NMR-MOUSE 278–279 PA–FTIR spectra 57–58, 62–63, 66–67 silica-filled 66–67 structure 401 styrene determination 62–63, 88, 89, 408–410 sulfur vulcanisation 338 thermal degradation 24–25 thermo-oxidative ageing 274–275

634

vinyl SBR 63–64, 336 see also carboxylated styrene– butadiene rubber Sulfonation ionomers 149–151, 198–199 IR spectra 138, 149–151, 198–199 maleated styrene/ethylene–butylene/ styrene 149–151 Sulfur vulcanisation accelerated 216–217, 327–333, 386– 387 accelerators 215, 322–323, 328–329, 330–333, 336, 506–507 activators 215, 323 cis-butadiene rubber 212–213, 333– 338 carbon-black-filled rubbers 336–337, 341–347, 582, 583 covulcanisation 269–271 dialkenylsulfides 217–219 diene rubbers 323–324 elastomers 209–210, 264 EPDM 208–209, 216–225, 238, 365, 386–387, 422 free radicals 213, 335 1 H NMR transverse magnetisation relaxation 273–274 mechanism 322–324, 345–346 model studies 209–210, 216–217 natural rubber 214, 323–333, 334, 582, 583 accelerated 327–333 unaccelerated 327 network chain density 333, 334, 344, 346–347, 367 NMR imaging 264, 273–274, 303– 304 optical spectroscopy 211, 214 pendent groups 323 polydiene rubbers 211–216 mechanism 214–216

cis-polyisoprene 330–333, 340, 343– 344 Raman spectroscopy 211–213, 214 Fourier transform Raman spectroscopy 217–223 SBR 338 silica-filled rubbers 347 sulfurating agents 323 see also vulcanisates Surface analysis 66–70 Swollen state chain structure 302, 303, 304, 306, 344, 346–347 cyclohexane solvent 266–267, 497– 498, 501 high-resolution NMR 443–446, 510 1 H NMR transverse magnetisation relaxation 354, 361–363, 378, 500– 506 2 H NMR 580–581 network density 502–510 NMR imaging 499–502 NR 506–508 polymer chain motion 494–496 rubbery materials 354, 491–518 self-diffusion 496–499 small molecule motion 496–500, 510, 580, 581

T Tacticity EPM rubber 90 IR spectra 90, 111 Telechelic chains 299, 423–424, 439 Temperature dependence chain structure 301–302 1 H NMR transverse magnetisation relaxation 356–357, 375, 376 log (shear) modulus 167–168 NMR imaging 276

129

Xe NMR spectroscopy 470, 471– 473, 476–478, 479 Temporary networks 310–312 Tenax absorption cartridges 3–4, 12, 16 Terminal groups 424–426, 436–437, 438 see also end-groups Termonomers, EPDM 90, 91 Tetramethyl thiuram disulfide (TMTD) 328, 334 TGA see (multi) hyphenated Thermogravimetric analysis (TGA) techniques; thermogravimetric analysis Thermal analysis 1 Thermal conductivity 51 Thermal decomposition carboxynitroso rubber 28 EPDM 23–24 pressure behaviour 26–27 sulfur vulcanisation 323 Thermal degradation ABS 21–22, 26 Aflas FA-150C rubber 21–22 carbon-black-filled SBR 24–25 IR spectra 113, 138, 156, 184, 185– 187 (multi) hyphenated TGA 17, 27–28 polyurethanes 194 PVC 184, 185–187 Thermal desorption 34–35 Thermal diffusion depth 52, 66–67 Thermal diffusivity 51–52, 54, 66–67 Thermal stability, (multi) hyphenated TGA 17 Thermobalances 3, 6, 13 Thermogravimetric analysis (TGA) 1 calibration 13–14 high pressure 4 multicomponent separation 12–13 quantitative analysis 18–19 sample requirements 12–13 see also (multi) hyphenated

635

Spectroscopy of Rubbers and Rubbery Materials thermogravimetric analysis (TGA) techniques Thermo-oxidative degradation 13 C NMR 348 network structure 347–348 NMR imaging 274–275 PVC 185–187 SBR 274–275 see also ageing; oxidation Thermoplastic elastomers (TPE) 2 H NMR 548–549, 585–588 IR spectra 188–199 recycling 188 Thermoplastic rubbers 168 Time-of-flight mass spectrometry (ToFMS) 10, 28, 36 Topology 353, 361, 378 TPE see thermoplastic elastomers Transmittance 82 Transparency 51 Triallylcyanurate (TAC) 229, 235, 502 Trimethylolpropane trimethacrylate (TMPTMA) 156–157 Two-dimensional NMR techniques 519– 520, 528–550 applications 530–549 double-quantum MAS spectroscopy 540–543 exchange spectroscopy 529, 530–535, 548–549 heteronuclear correlation experiments 543–547 2 H NMR 548–549 multi-quantum spectroscopy 538–543 NOESY–MAS spectroscopy 535–538 separation spectroscopy 529 Tyres, NMR imaging 267–269

U Urey–Bradley field theory 126

636

V Vibrational frequencies 78–80, 79, 170– 172 Vinyl acetate 60–61 Vinyl silanes, grafting onto polyethylene 174, 175, 176 Vinyl styrene–butadiene rubber 13 C NMR 336 PA–FTIR spectrum 63–64 Viscoelastic materials 1 H NMR transverse magnetisation relaxation 365–366, 383–384, 526– 527 2 H NMR 522, 527 ionic 383–384 multidimensional NMR characterisation 519–555 property temperature dependence 167–168 Viscous melts 167 Volatiles, residual 26 Volume-average material properties 360 Vrentas–Duda equation 497 Vulcanisates 13 C NMR 322, 324–333, 333–338, 338–339 high-resolution solid state NMR 321– 352 rubbers 209–210 oil-extended 366 Vulcanisation see high-pressure Vulcanisation; peroxide-curing; sulfur Vulcanisation

W Wavelength 77–78 Weathering, IR spectra 138, 153 Welding, crosslinked polyethylene pipes 178, 179 Wide-line separation experiments (WISE)

543–547 Williams–Landel–Ferry (WLF) equation 494, 497–498, 586

X 129

Xe NMR spectroscopy chemical shift 457–458, 469–470, 476–480 elastomers in blends and composites 457–489 EPDM 458, 463–464, 474–480, 482– 486 EPM 465–474, 482 exchange model 479–480, 484 gas phase 458–459 PBT/PTMO block copolymer 486–487 polyisoprene 485–486 polymers 459–461 pulsed field gradient 462–463, 480– 482, 483–484, 485 self-diffusion coefficients 461, 463, 464, 480–487

diffusion time 485–487 temperature dependence 470, 471– 473, 476–478, 479 XNBR see carboxylated nitrile rubber XSBR see carboxylated styrene–butadiene rubber

Y Young’s modulus 460

Z Ziegler–Natta catalysis 168, 207 Zinc oxide blends, IR spectra 151–153, 154–155, 195–198 sulfur vulcanisation 215, 340 Zinc stearate bloom 66–68 IR spectra 198–199 PVC thermal stability 187

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Spectroscopy of Rubber and Rubbery Materials

Spectroscopy of Rubber and Rubbery Materials

Thermal Analysis of Rubbers and Rubbery Materials

Vibrational Spectroscopy of Biological and Polymeric Materials

Infrared and Raman Spectroscopy of Biological Materials

Infrared and Raman Spectroscopy of Biological Materials

Vibrational Spectroscopy of Biological and Polymeric Materials

Rubber and Rubber Balloons

Analysis of rubber and rubber-like polymers

Rubber

Rubber

Rubber

Rubber

Rubber

Rubber

Scanning Probe Microscopy of Functional Materials: Nanoscale Imaging and Spectroscopy

Blends of Natural Rubber

Physical Testing of Rubber

The Romance Of Rubber

Handbook of Rubber Bonding

Mixing of Rubber

Plastics, Rubber and Health

Rubber Soul

Rubber Soul

Burning Rubber

Burning Rubber

Spectroscopy

Spectroscopy

Rubber Basics

Rubber Recycling

Determination of Additives in Polymers and Rubber

Spectroscopy of Rubber and Rubbery Materials