Handbook of Engineering Polymeric Materials

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HANDBOOK OF ENGINEERING POLYMERIC MATERIALS

edited by

Nicholas P. Cheremisinoff, Ph.D.

m MARCEL

D E K K E R

MARCEL DEKKER,

INC.

N E WYORK BASEL HONGKONG

Library of Congress Cataloging-in-Publication Data

Handbook of engineering polymeric materials / edited by Nicholas P. Cheremisinoff. cm. p. Includes index. ISBN 0-8247-9799-X (hc:a l k . paper) 1. Polymers. 2. Polymeric composites.3. Elastomers. I. Cheremisinoff, Nicholas P. TA455.P58H36 1997 668.94~21 97-20896 CIP

The publisher offers discounts on this book when ordered in bulk quantities. For more information, write to Special SalesProfessional Marketing at the addressbelow. This book is printed on acid-free paper. Copyright 0 1997 by Marcel Dekker, Inc. All Rights Reserved. Neither this book nor any partmay be reproduced or transmitted in any form or by any means, electronicor mechanical, including photocopying, microfilming, and recording, or by any information storage and retrieval system, without permission in writing from the publisher. Marcel Dekker, Inc. 270 Madison Avenue, New York, New York 10016 Current printing (last digit): I O 9 8 7 6 5 4 3 2 1

OF AMERICA PRINTED IN THE UNITED STATES

Preface

Polymers and, in particular, elastomers have evolved into mature products over the last thirty with years,wellestablished applications, mostof which center around the consumer. Despite these materials being mature, and market applications well established, this is by no means a stagnant industry or subject. Driven by fierce competition, product quality improvements, and new applications in such market segments as health care, automotive parts, construction and building materials, mechanical goods, and an enormous range of composites of elastomers, plastics blends, and polywith specialty or niche-type applications, suppliers and manufacturers meric alloying materials continueto strive fornew products and variations of feedstock polymers. This volume provides a compendium of some of the latest technological advancements in product applications, new elastomers, blends, alloys, and functionalized materials. Discussionsare also included on some of the more conventional polymers to provide a balance to the discussions of alternative materials for end-use applications. As already noted, although the elastomers and plastics industries may be mature, they arestill very active. Products and raw materialsare constantly evolving; new ones are introduced while others are dropped. Thermoplastic rubbers began to emerge in the late 1980s and have now made substantial market penetrations into automotive applications, mechanical goods, construction materials, and other areas. Perhaps more than plastics, it is the rubbers or elastomers that continue to show the greatest areas of product innovation because of their versatility in properties, but also their ability for compatibility with certain plastic materials. In general, the demand for synthetic rubber is strong throughout the world, particularly iR North America. We should not forget, however, that natural rubber continues to account for25 to 30% of total elastomer demand, and will likely remain strong in the automotive market sector as the move to steel radial truck tires from bias-ply tires continues. This volume concentrates on the plastic-elastomeric types of blends and alloys and examines both new and conventional products. Emphasis in discussions is given to product characterizationand performance attributes, as well as the structural properties of these materials. In this regard the volume is probably most useful to the product development specialist and the applications engineer. The book represents the efforts of a large number of experts from both the industrial and academic communities. Their in preparing efforts contributions to this volumeare to be noted and I express heartfelt gratitude for theirtime and effort. A special thanks is extended to thepublisher for the fine productionof this volume.

Nicholas P . Cheremisinoff, Ph.D.

iii


Contents

Preface iii Contributors ix 1. Artificial Neural Networks as a Semi-EmpiricalModeling Tool for Physical Property Predictions in PolymerScience 1 Ramesh Keshavaraj, Richard William Tock, and RaghuS . Narayan 2. New Generation High Performance Polymers by Displacement Polymerization 35 Pradip Kumar Dutta, Susanta Banerjee, and Sukumar Maiti 3.

AcrylamidePolymers Valery F. Kurenkov

4.

Transparent Polyolefins 73 Yu. P . Losev

5.

Polyolefin Stabilizers with Intramolecular Synergism 77 Yu. P . Losev

61

6. High-Temperature Stabilization of Polyolefins 81 Yu. P . Losev 93

7.

Poly(ma1ic Acid) from Natural Sources Eggehard Holler

8.

Stabilization of Polyolefins Sahar AI-Malaika

9.

Gamma Radiation Induced Preparation of Polyelectrolytes and Its Use for Treatment of WasteWater 119 T. Siyam

105

10. Polyvinylchloride (PVC)/Thermoplastic Polyurethane (TPU) Polymeric Blends Cheng-Di Huang and Jiong Chen 11. Mechanical Properties of Ionomers and Ionomer Blends Masanori Hara and J . A . Sauer

145

12. Metallocene-Based Polyolefins: ProductCharacteristics Virendra Kumar Gupta

155

V

137

vi

Contents

13. Preparations and Properties of Porous Poly(viny1 alcohol)-Poly(viny1 acetate) Composites

Sadao Hayashi

14. Hydrophobization of Polyanionic Polymers to Achieve Higher

Biological Activity

Naoto Oku, Raphael M. Ottenbrite, and Yasuo Suda

167

179

15. Uniform Latex Particles 189 Ali Tuncel and Huseyin Cicek 16. Reaction Mechanism of Vinyl Polymerization with Amine in Redox and Photo-Induced Charge-TransferInitiation Systems 227

Xin-De Feng, Kun-Yuan Qiu, and Wei-Xiao Cao

17. Photoinitiation of Free Radical Polymerization by Organometallic Compounds

Salah M. Aliwi

243

18. Chemical Modification of Polystyrenes in the Presence of Cationic Catalysis and Their IndustrialApplications

259

Ramazan Mirzaoglu, Refiga Kurbanova, and Mustafa Ersoz 19. Performance of Polyethylenes in Relation to Their Molecular Structure

Ajit B. Mathur and Ishwar Singh Bhardwaj

20. The Crystallization of Polyethylene Under High Pressure

Rui Huang, Xiongwei Zhang,and Guangxian Li

21. 22.

Structure,Stability,andDegradation

Mohammad Kazim Naqvi

295

of PVC 317

Solution State of Metal Complex Calixarenes and Polymeric Calixarenes

Mustafa Yilmaz

23. Thermodynamic Opportunity Due to the Degradation Reaction Initiation

25. 26. 27.

Polymerization by Ylides 373 A. K. Srivastava CharacterizationandProperties of Thermotropic Polybibenzoates 383

Ernesto Ptrez, Jost M.Peretia, Rosario Benavente, and Antonio Bello Polymeric UltravioletStabilizersforThermoplastics

Jayant S. Parmar and Raj P. Singh

399

High-Performance and Functional Materials from Natural Monomers

Chennakkattu Krishna Sadasivan Pillai

28. Ionic Thermoplastic Elastomer Based on Maleated EPDM Rubber Santanu Datta and S. K . De 29.

339 351

E. F. Vainstein and G . E. Zaikov

24.

277

Elastomeric Compound-The

and Polymers 411 441

Importance of Consistency in Economic Component Manufacture

Robert Bond

30. Advancement in Reactively Processed Polymer Blends 465

Mohini M. Sain and Claude Daneault

31. GraftCopolymerization Onto Natural and Synthetic Polymers

Inderjeet Kaur and Bhupendra Nath Misra

32. Methods,Characterization,andApplications

of Grafting 501 Elsayed Mohamed Abdel-Baty and Eman Mohamed El-Nesr

481

451

vii

Contents 33.

Surface Properties of Modified Polymers 517

Jayasree Konar, Bhola Nath Avasthi, and Anil K . Bhowmick 34. Grafting of Cellulose 529 Abd-Alla M. A. Nada and Mohamed Adel Yousef 35.

Graft Copolymerization of Vinyl Monomers Onto Macromolecules Having Active Pendant Group via Ceric Ion Redox or Photo-Induced Charge-Transfer Initiation 541

Kun-Yuan Qiu and Xin-De Feng 36.

Modification of Properties of Nitrile Rubber

555

Susmita Bhattacharjee, Anil K. Bhowmick,and Bhola Nath Avasthi 37.

Industrial Perspective of Cellulosics in Thermoplastic Composites 577

Debesh Maldas 38.

Liquid Crystalline Polymer Composite-Preparation and Properties 585

Yongsok Seo 39.

Self-organization of Core-Shell Type Polymer Microspheres and Applications to PolymerAlloys 601

Reiko Saito and Koji Ishizu 40.

Flow Behavior of Polymer Blends as Affected by Interchain Crosslinking 61 1 Chapal K. Das

41.

Tough Composites Based on Premixed PPLCP Blends 623

Markku T. Heino and Tommi P . Vainio 42.

Compatibilization of Thermoplastic Elastomer Blends 633

R . Asaletha, Zacharia Oommen, and Sabu Thomas 43.

Polymer Blends and Alloys 649

B. S. Kaith, A. S . Singha, and Sunil

44. Reactive Compatibilization of Immiscible Polymer Blends Snooppy George, Josephine George, and Sabu Thomas 45.

667

Polymer Blends Containing Thermotropic Liquid Crystalline Polymer

685

Xiao-Su Yi 46.

EnergeticComposites

705

M. B. Khan 47.

Macrointermediates for Block and Graft Copolymers

725

Baki Hazer 48.

Azo Initiators as Transformation Agents for Block Copolymer Synthesis 735

Yusuf Yagci and Zvo Reetz 49.

Block Copolymers Derived from Macroinitiators: Recent Advances of Synthesis, Properties, andApplications 755

Akira Ueda and Susumu Nagai 50. Poloxamers 765

Iain F. Paterson, Babur 2. Chowdhry, and Stephen A. Leharne

51.

Recent Advances in Cationic Resins

775

Metwally Shafik Metwally 52. Natural Fiber Reinforced Plastics A. K. Bledzki and J . Gassan

787

viii

53.

Contents

ShortFiber-ReinforcedPlasticComposites Jayamol George and Sabu Thomas

81 1

54. Peculiarities of the Fine Structure of PET Fibers and the Relationship to Their Basic Physical Properties 839 Grzegorz Urbahczyk and Andrzej Jeziorny

Index 859

Contributors

Elsayed Mohamed Abdel-Bary Department of Chemistry, Mansoura University, Mansoura, Egypt Sahar Al-Malaika Chemical Engineering and Applied Chemistry, Aston University, Birmingham, England Salah M. Aliwi Department of Chemistry, College of Science, Mustansiriya University, Baghdad, Iraq R. Asaletha*

Rubber Research Institute of India, Kerala, India

Bhola Nath Avasthi Department of Chemistry, Indian Institute of Technology, Kharagpur, India Susanta Banerjee Department of Synthetic Chemistry (Polymer Group), Defence Establishment, Gwalior, India Antonio Bello

Research 8z Development

Instituto de Ciencia y Tecnologia de Polimeros (CSIC), Madrid, Spain

Rosario Benavente

Instituto de Ciencia y Tecnologfa de Polimeros (CSIC), Madrid, Spain

Ishwar Singh Bhardwaj Susmita Bhattacharjee

R. S. Petrochemicals Ltd., Baroda, India FB-Physikalische Chemie-Polymere, Philipps Universitat, Marburg, Germany

Ani1 K. Bhowmick Rubber Technology Center, Indian Institute of Technology, Kharagpur, India

A. K. Bledzki Institute of Materials Technology, Plastic and Kassel, Germany Robert Bond

Recycling Technology, University of Kassel,

Chicago Rawhide (Americas), Elgin, Illinois

Wei-Xiao Cao Institute of Polymer Science, Peking University, Beijing, China Jiong Chen Department of Technical Development, Shanghai Medical Equipment Research Institute, Shanghai, China Babur Z. Chowdhry School of Chemical and Life Sciences, University of Greenwich, London, England Huseyin Cisek Chemical Engineering Department, Hacettepe University, Ankara, Turkey Claude Daneault Pulp and Paper Research Center, University of Quebec, Trois-Rivieres, Quebec, Canada Chapal K. Das Materials Science Centre, Indian Institute of Technology, Kharagpur, India Santanu Datta Rubber Technology Centre, Indian Institute of Technology, Kharagpur, India

S. K. De Rubber Technology Centre, Indian Institute of Technology, Kharagpur, India

* Current uflliution: Mahatma Gandhi University, Kerala, India ix

Contributors

X

Pradip Kumar Dutta Department of Applied Chemistry, Shri G. S. Institute of Technology & Science, Indore, India Eman Mohamed El-Nesr

National Center for Radiation Research and Technology, Cairo, Egypt

Mustafa Ersoz Department of Chemistry, Selsuk University, Konya, Turkey Xin-De Feng

Institute of Polymer Science, Peking University, Beijing, China

J. G u a n Institute of Materials Technology, Plasticand Recycling Technology, University of Kassel, Kassel, Germany Jayamol George School of Chemical Sciences, Mahatma Gandhi University, Kerala, India Snooppy George School of Chemical Sciences, Mahatma Gandhi University, Kerala, India Josephine George

School of Chemical Sciences, Mahatma Gandhi University, Kerala, India

Virendra Kumar Gupta Research Centre, Indian Petrochemicals Corporation Ltd., Gujarat, India Masanori Hara Department of Chemical and Biochemical Engineering, Rutgers University, Piscataway, New Jersey Sadao Hayashi Faculty of Textile Science and Technology, Shinshu University, Ueda, Japan Baki Hazer Department of Chemistry, Zonguldak Karaelmas University, Zonguldak, Turkey Markku T.Heino* Department of Chemical Engineering, Helsinki University of Technology, Espoo, Finland Eggehard Holler Institute for Biophysics and Physical Biochemistry, Department of Biology and Preclinics, University of Regensburg, Regensburg, Germany Cheng-Di Huang Department of Technical Development, Shanghai Medical Equipment Research Institute, Shanghai, China Rui Huang Department of Plastics Engineering, Sichuan Union University, Sichuan, China Koji Ishizu Department of Polymer Science, Tokyo Institute of Technology, Tokyo, Japan Andrzej Jeziorny Institute of Fibre Physics and Textile Finishing, Technical University of Lodt, Lodi,Poland

B. S. Kaith Department of Applied Sciences and Humanities, Regional Engineering College, Hamirpur, India Inderjeet Kaur Department of Chemistry, Himachal Pradesh University, Shimla, India Ramesh Keshavarajt Department of Chemical Engineering, Texas Tech University, Lubbock, Texas M. B. Khan CHEMTEC & Prime Glass, Jhelum, Pakistan Jayasree Konar Department of Chemistry, Indian Institute of Technology, Kharagpur, India Refiga Kurbanova Department of Chemistry, SelGuk University, Konya, Turkey Valery F. Kurenkov Kazan State Technological University, Kazan, Russia Stephen A. Leharne

School of Earth andEnvironmental Sciences, Universityof Greenwich, London, England

Guangxian Li Department of Plastics Engineering, Sichuan Union University, Sichuan, China Yu. P. Losev Department of Chemistry, Belarussian State University, Minsk, Belarus Sukumar Maiti Materials Science Centre, Indian Institute of Technology, Kharagpur, India Current afiliations:

* Nokia Cables Ltd., Espoo, Finland t Milliken & Company, LaGrange, Georgia

xi

Contributors

Debesh Maldas* Pulp and Paper Research Center, University of Quebec, Trois-Rivieres, Quebec, Canada Ajit B. Mathur Research Centre, Indian Petrochemicals Corporation Ltd., Baroda, India Metwally Shafik Metwally Department of Chemistry, Al-Azhar University, Cairo, Egypt Ramazan Mirzaoglu Department of Chemistry, Selguk University, Konya, Turkey Bhupendra Nath Misra

Department of Chemistry, Himachal Pradesh University, Shimla, India

AM-Alla M. A. Nada Cellulose & Paper Department, National Research Centre, Cairo, Egypt Susumu Nagai Plastics Technical Society, Osaka, Japan Mohammad Kazim Naqvi PVC Group, Researchand Technology Support, Saudi Basic Industries Corporation (SABIC), Riyadh, Saudi Arabia Raghu S. Narayan Department of Chemical Engineering, Texas Tech University, Lubbock, Texas Naoto Oku Department of Radiobiochemistry, School of Pharmaceutical Sciences, University of Shizuoka, Shizuoka, Japan Zacharia Oommen C.M.S. College, Mahatma Gandhi University, Kerala, India Raphael M. Ottenbrite Department of Chemistry, Virginia Commonwealth University, Richmond, Virginia Jayant S. Parmar Department of Chemistry, Sardar Patel University, Gujarat, India Iain

F. Paterson School of Earth and Environmental Sciences, University of Greenwich, London, England Instituto de Ciencia y Tecnologia de Polimeros (CSIC), Madrid, Spain

J o d M. Pereiia

Emesto P6rez Instituto de Ciencia y Tecnologia de Polimeros (CSIC), Madrid, Spain Chennakkattu Krishna Sadasivan Pillai Polymer Division, Regional Research Laboratory(CSIR), Thiruvananthapuram, India Kun-Yuan Qiu Institute of Polymer Science, Peking University, Beijing, China Ivo Reetz Organic Chemistry Department, Istanbul Technical University, Istanbul, Turkey Mohini M. Sain

Pulp and Paper Research Center, University of Quebec, Trois-Rivieres, Quebec, Canada

Reiko Saito Department of Polymer Chemistry, Tokyo Institute of Technology, Tokyo, Japan J. A. Sauer Department of ChemicalandBiochemicalEngineering,RutgersUniversity,Piscataway,

NewJersey

Yongsok Se0 Polymer Processing Laboratory, Korea Institute of Science and Technology, Seoul, Korea Raj P. Singh

Polymer Chemistry Division, National Chemical Laboratory, Pune, India

A. S. Singha Department of Applied Sciences and Humanities,Regional Engineering College,Hamirpur, India T. Siyam Nuclear Chemistry Department, Hot Laboratory Centre, Atomic Energy Authority, Cairo, Egypt

A. K. Srivastava Chemistry Department, Harcourt Butler Technological Institute, Kanpur, India Yasuo Suda Department of Chemistry, Osaka University, Osaka, Japan Sunil Department of Applied Sciences, Regional Engineering College, Hamirpur, India Sabu Thomas School of Chemical Sciences, Mahatma Gandhi University, Kerala, India Richard William Tock Department of Chemical Engineering, Texas Tech University, Lubbock, Texas

* Current

afJiliation: Michigan State University, East Lansing, Michigan

xii

Contributors

Ali Tuncel Chemical Engineering Department, Hacettepe University, Ankara, Turkey Akira Ueda Osaka Municipal Technical Research Institute, Osaka, Japan

Grzegorz Urbaliczyk Institute of Fibre Physics and Textile Finishing, Technical University of Lodt, Lodt, Poland Tommi P. Vainio* Department of Chemical Engineering, Helsinki University of Technology, Espoo, Finland E. F. Vainstein Department of Kinetics and Thermodynamics of Cooperative Processes, N. M. Emanuel Institute of Biochemical Physics, Russian Academyof Sciences, Moscow, Russia Yusuf Yagci Organic Chemistry Department, Istanbul Technical University, Istanbul, Turkey Xiao-Su Yi

Institute of Polymers and Processing, Zhejiang University, Hangzhou, China

Mustafa Yilmaz Department of Chemistry, SelGuk University, Konya, Turkey Mohamed Adel Yousef Department of Chemistry, Helwan University, Cairo, Egypt

G . E. Zaikov Department of Chemical and Biological Kinetics, N. M. Emanuel Institute of Biochemical Physics, Russian Academy of Sciences, Moscow, Russia Xiongwei Zhang Department of Plastics Engineering, Sichuan Union University, Sichuan, China

* Current uffiliarion:Valmet Corporation, Jyviiskyla, Finland

Artificial Neural Networks as a Semi-Empirical Modeling Tool for Physical Property Predictions in Polymer Science Ramesh Keshavaraj,* Richard William Tock, and Raghu S. Narayan Texas Tech University, Lubbock, Texas

1.

INTRODUCTION

pology of a directed graph, which can carry out information processing, by means of its outputs, on either on-off or continuous inputs-HechtNielsen

Recently, a new approach called artificial neural networks (ANNs) is assisting engineers and scientists in their assessment of “fuzzy information.” Polymer scientists often face a situation where the rules governing the particular system are unknown or difficult to use. It also frequently becomes an arduous task to develop functional forms/empiricalequations to describe a phenomena. Most of these complexities can be overcome with an ANN approach because of its ability to build an internal model based solely on the exposure in a training environment. Fault tolerance of ANNs has been found to be veryadvantageousin physical property predictions of polymers. This chapter presents a few such cases where the authors have successfully implemented an ANN-based approach for purpose of empirical modeling. These are not exhaustive by any means. Typical applications of neural networks have been largely made in the areas of pattem recognition, signal processing, speech recognition, written character recognition, time series analysis, complex dynamic systems, and process control [l-141. It is important to note that allof these applicationsinvolveparallelinformation transformationbetween various processing elements connected through different architectures. The specific application of ANNs to physical property prediction, especially with respect to polymers, has been rather limited.

II. ANN AS A MODELING TOOL The computational paradigmin an ANN is based on an idealized modelof a biological unit calleda neuron. The unique characteristics of this ANN modelare the inputs of signal from stimulus in a training environment. It is important to note that each neuron works independently of the other neurons. The specific characteristics of ANN models that attract industrial application are: 1. 2. 3. 4. 5.

Learning from examples Continued learning while inoperation Ability to distinguish noise in experimentaldata Ability to generalize Self-organization 6. High speed in conjunction with complex parallel processing 7 . Error tolerance These capabilities of ANNs make them a unique tool for a large number of industrial applications. In this chapter, the authors demonstrate, with case studies, the advantages of using this approach to physical property predictions in polymer science.

A artificial neural network is to a great extent, a parallel-processing dynamic system, with the to-

111.

FUNDAMENTALS OF ANNs

To understand why and how artificial neural networks work as they do, it is helpfulto study some of the funda-

*Current afJiliation: Milliken & Company, LaGrange, Georgia.

l

Keshavaraj et al.

2

mental characteristics of biological neural networks. A biological neuron is a primary processing unit.In a biological neuron, dendrites carry information impulses toward the neuron andcan be considered as an input. Similarly, an axon carries information impulses away from the neuron and can beconsidered as an output. Hence, artificial neuralnetworks are constructed of mathematical processingelements, also known as neurons, which are connected via informationchannels called interconnections/synapse, as in a biological neuron. Synapses in a network serve as pathways of varying strength for transmission of information. The neurons operate collectively and simultaneously and are configured in regular architecture. They “learn” by extracting preexisting information from the data that describe the relationship between the inputs and the outputs. Hence, in the learning process, the network actually acquires knowledge or information fromthe input environment. As a result of the interrelationships, the network assimilates trends and relationships that can be recalled later. Neural net-

works can identify and learn correlative patterns between sets of input data and correspondingtarget values. Once trained, neural networks can be used to forecast the outputs expected for new levels of input variables. Each neuron can have multiple inputs, but only one output. Each output may, however, branch out as an input to other neurons. Neural networks that are capable of handling complex and nonlinear problems can process information rapidly andhelp reduce the engineering effort required in developing highlycomputation-intensive models. Neural networks also come ina variety of types, and each has its distinct architectural differences and uses [15,16]. The structure of a neural network forms the basis for information storage and governs the learning process. The type of neural network used in this work is known as a feed-forward network; the information flows only inthe forward direction, i.e., from input to output in the testing mode. A general structure of a feed-forward network is shown in Fig. 1. Connections are made be-

NETWORK OUTPUT

NETwoRgINPuT

Figure 1 A general structure of a feed-forward neural network.

3

Artificial Neural Networks as a Modeling Tool

tween neurons of adjacent layers: a neuron isconnected so that it receives signals fromeach neuron inthe immediate precedinglayer and transmits signals to each neuron inthe immediate succeedinglayer. Neural networks that are organized in layers typically consist of at least three layers: an input layer, one or more hiddenlayers, and an output layer. The input and output layers serve as interfaces that perform the appropriate scaling relationship betweenthe actual and the network data. Hidden layers are so termed because their neurons are hidden from the actual data; the connections provide the means for information flow.Each connection has an associated synaptic weight factor, wi, expressed by a numerical value that can be adjusted. The weight is an indication of the connection strength between any two neurons. The neurons in both the hidden and output layers perform summing and nonlinear mapping functions. The functions carried out by each neuron are illustrated in Fig. 2. Each neuron occupies a particular position in a feed-forward network andaccepts inputs only from the neurons in the preceding layer and sends its outputs to other neurons in the succeeding layer. The inputs from other nodes are first weighted and then summed. This summing of the weighted inputs is carried out by a processor within the neuron. The sum that is obtained is called the activation of the neuron. Each activated neu-

ron performs three primary functions: (1) receives signals from other neurons; (2) it sums their signals; and (3) it transforms the sum. For example, if the output from the ith neuron withpattern p is designated as xi.p, then the input to the jth neuron from the ith neuron is xiSpwij.Summing the weighted inputs to the jth neuron can be represented as: i

where t?j is a bias term andW B is~ the weight of the connection fromthe bias neuronto thejth neuron. The bias neuron is provided to supply an invariant output to each neuron in the hidden layer. If the weighted sum of the scaled inputs into the hidden layers exceeds the weight of the bias hidden layer neuron, the neuron becomes activated or excited. This activation can either be positive, zero, or negative, because the synaptic weightings and the inputs can beeither positive or negative. A bias of 1 was used in this study. Hence, any weightedinput that makes a positive contribution to activation represents a triggering or tendency to turn the neuron on. An input making a negative contribution represents an inhibition, whichtends to turn the neuron off. After summing its inputs to determine its activation, the summed total is then modified by a mapping function, also known as a transfer/threshold function. A commonlyused transfer function is the “sigmoid,” which is expressed as:

+

Output signal ‘S’ to other nodes

Transfer function

(sigmoidal)

neuron

Activation (summation)

-

ww l j 2wj 3 j

x1

x2

synapic

x3

Inputs from other nodes

Figure 2 Schematic of typical functions carried outby a

neuron.

A sigmoid (S-shaped) is a continuous function that has a derivative at all points and is a monotonically increasing function. Here Si,pis the transformed output asymptotic to 0 S Si,p I1 and ui,pis the summed total of the inputs (- m S ui,pI+ W) for pattern p . Hence, when the neural network ispresented with a set of input data, each neuron sumsup all the inputs modified bythe corresponding connection weights and appliesthe transfer function to the summed total. This process is repeated until the network outputs are obtained.

W.

NEURAL NETWORKARCHITECTURE

The number of neurons to be used in the inputloutput layer are based on the number of inputloutput variables to be considered in the model. However, no algorithms are available for selecting a network structure or the number of hidden nodes. Zurada[l61 has discussed several heuristic based techniques for this purpose. One hidden layeris more than sufficient for most problems. The number of neurons in the hidden layer neuron was selected by a trial-and-error procedure by monitoring the sum-of-squared error progression of the validation data set used during training. Details about this proce-

4

Keshavaraj et al.

dure are discussed in the following paragraphs. A feedforward type of neural network architecture has been used in all the case studies considered in this chapter. The training algorithm usedhere is based onthe familiar Fletcher algorithm for nonlinear optimization used in many thermodynamic property estimations. V. TYPES OF ANNs

Neural networks can be broadly classified based on their network architecture as feed-forward andfeed-back networks, as shown in Fig. 3. In brief, if a neuron's output is never dependent on the output of the subsequent neurons, the network is said to be feed forward. Input signals go onlyone way, and the outputs are dependent on only the signals coming in from other neurons. Thus, there are no loops in the system. When dealing withthe various types of ANNs, two primary aspects, namely, the architecture and the types of computationsto be per-

FEED-BACK NEURAL NETWORKS

:

...

formed, have to be understood before their implementation. Neural networks canalso be classified bytheir neuron transfer function, which typically are either linear or nonlinear models. The earliest modelsusedlinear transfer functions wherein the output values were continuous. Linear functions are not very useful for many applications because most problemsare too complex to be manipulated by simple multiplication. In a nonlinear model, the output of the neuron is a nonlinear function of the sum of the inputs. The output of a nonlinear neuron can have a very complicated relationship with the activation value. A.Feed-BackNetworks In a feed-back network, the output signals of neurons directly feed back into neurons in the same or preceding layers, as shown in Fig. 3. Feedback refers to the way in which neurons are connected. Feed-back networks send signals backas inputs to other neurons. This is not same as back-propagation, which describes a training method. With back-propagation, the error at the final output layer is usedfor correction. Feed-back networks do not use back-propagation for training. Back-propagation networks are not feed-back models. Feed-back models can be constructed and trained. In a constructed model, the weight matrix iscreated by adding the output product of every input pattern vector with itself or with an associated input. After construction, a partial or inaccurate input pattern can be presented to the network and, after a time, the network converges to one of the original inputpatterns. Hopfield and BAM are two well-known constructed feed-back models.

B. Feed-ForwardNetworks The second main category of neural networks is the feedforward type. In this type of network, the signals go in only one direction; there are no loops in the system as showninFig. 3. The earliest neural network models were linear feed forward. In 1972, two simultaneousarticles independently proposed the same model for an associative memory, the linear associator. J. A. Anderson [17], neurophysiologist, and Teuvo Kohonen[18], an electrical engineer, wereunaware of each other's work. Today, the most commonly used neural networks are nonlinear feed-forward models. Current feed-forward network architectures work better than the current feed-back architectures for a number of reasons. First, the capacity of feed-back net&ownwithout o Bios works is unimpressive. Secondly, in the running mode, feed-forward modelsare faster,since they need to make Figure 3 Feed-back and feed-forward artificial neural net- one pass through the system to find a solution. In conworks. trast, feed-back networks must cycle repetitively until FEED-FORWARD NEURAL NETWORK


the neuron outputs stop changing. This cycling can typically require anywhere from 3 to lo00 cycles. Computations in neuralnetworks can also be broadly classifiedinto two other categories, namely, direct throughput andrelaxation. Direct throughputis specific to feed-forward neural networks. In this case, the selected network trainingdata are trained so that a certain set of input patterns give rise to desired outputs. Relaxation, to the contrary, is an iterative convergence to a fixed point, which is achieved by feed-back loops constantly feedingthe network output into the neurons in the preceding layers. These two modified networks can be used advantageously in optimization problems. Also, relaxation networks, when correctly trained, can also be used as content addressable memory.

VI. LEARNING IN NEURALNETWORKS

5

learning process is to adjust the weights inorder tominimize the error between the network predicted output and the desired output. The output from each neuron i is Si,p, as shown in Eq. (2).

A. NonlinearOptimizationTraining Routine A faster training process is to search for the weights with the help of a optimization routine that minimizes the same objective function. The learning rule used in this work is common to a standard nonlinear optimization or least squares technique. The entire set of weights are adjusted at once instead of adjusting them sequentially from the output to the input layers. The weight adjustment is done at the end of each exposure of the entire training set to the network, and the sum of squares of all errors for all patterns is usedas the objective function for the optimization problem. A nonlinear optimization routine based on the Levenberg-Marquardt method [21] is used for solving the nonlinear least-squares problem. The optimization problem can be defined if the model to be fitted to the data is written as follows:

Once the network architecture is selected and the characteristics of the neurons and the initial weightsare specified, the network hasto be taught to associate new patterns and new functional dependencies. Learning corresponds to adjustments of the weights in order to F(Y) = f(a1, W , ,a m ; obtain satisfactory input-output mapping. Since neural p19 &, P k ) = f(a, p) (3) networks do not use “a priori” information about the process to be modeled, learning must come from trainingwhere al,a2, . , a,,,are independent variables, pl, in the exposed environment. Hence, the selection of p2, . . . ,P k are the population valuesof k parameter and training data in any problem becomes very important F ( y ) is the expected value of the independent variable y. for neural networks. The weights are learned through Then the data points can be denoted by: experience, using an update rule to change the synaptic ( Y i , X l i , Xzi, . . ,X,,,,-) i = 1, 2, . . . , n (4) weight, W“. Learning can be broadly dividedinto three categorThe problem isto compute those estimates of the paramies: supervised, unsupervised, and reinforced learning. eter that will minimizethe following objective function: 1. Supervisedlearning-Atraining data set containing a set of inputs and target outputs are provided to the selected network architecture. The synaptic weights are adjusted so as to miniwhere is the value of y predicted by the model at the mize the error between the desired and the netith data point. The parameters to be determinedis work predicted outputs for each input pattern. the strength of the connections, i.e., the weights, wt. A trained network can then be validated witha If the (m x n) Jacobian is defined by Jii = ari/axj(ri is secondindependent data set over the same the residual Yi - pi) then each iteration can be written range as the training data set. as follows: 2. Unsupervisedlearning-In this type the netYk = Yk 6k (6) work is able to discover statistical regularities in its input space and automatically develops where 6 is the solution of the set of linear equations of different modesof behavior to represent differthe following form: ent types of inputs. ( A AI)S = - V (7) 3. Reinforced-This type of network receives a globalreward/penaltysignalduringtraining. where A = JTJ and v = JTr are evaluated at Yk, and Usually weights are changed so as to develop where h is an adjustableparameter that is used to control an inpudoutput behavior that maximizes the the iteration. Moredetails of this Levenberg-Marquardt probability of receiving a reward, i.e., minimize method can be foundelsewhere [21]. A flow diagramof that of receiving a penalty. this original Marquardt algorithm is given in Fig.4. This algorithmshares with the gradient methods the Several different learning rules have been proposed ability to converge from an initial guess that may be by various researchers [15,19,20], but the aim of every

-..

.

+

+

9

6

Keshavaraj et al.

I

,

Figure 4 OriginalLevenberg-Marquardtalgorithm.

outside the range of convergence with other methods. It shares with the Taylor series method the ability to close in onthe converged values rapidly after the vicinity of the converged values have beenreached. 1. Modifications of Levenberg-Marquardt Method: Fletcher’s Modification A number of modifications to eliminate some less favorable aspects of the Levenberg-Marquardt method were considered by Fletcher. For instance, the arbitrary initial choice of the adjustable parameter A , if poor, can cause an excessive number of evaluations of (b,the sum of squared error, before a realistic value is obtained. This is especially noticeableif v, i.e., JTR(x), is chosen to be small, i.e., v = 2. Another disadvantage of the method is that the reduction of A to A h at the start of each iteration may also cause excessive evaluations, especially when v is chosen to be large, i.e., v = 10. The effect of this is that the average number of evaluations

of (b per iteration may be about 2, which isunnecessarily inefficient. A further disadvantage of the method isthat the tests 4’ I(b or even (b’ < (b precludes a proof of convergence being made. Finally, when solving problems in which R = 0 at the solution, it is possible to achieve a quadratic rate of convergence with the Gauss-Newton method, but only a superfine rate with Marquardt scheme.All of these drawbacks of the original Marquardt method were overcome using Fletcher’s modification [21]. Details about this implementationare given in Fig.5. Where R is the ratio of the actual reductiodpredicted reduction. The motivation for Fletcher’s modification was that if the ratio R is near 1, then the adjustableparameter A has to be reduced, and if the ratio is near to or less than zero, than A has to be increased. Fletcher suggested an arbitrary choice of constants p and u such that 0 C p < u < 1, and a reduction of A if R < p , and anincrease of A if R u.Various experiments were performedby Fletcher to study the effect of these

7


r Solve (A+l.9 6 = -v for 6 Evdtutc

Cdcnlate a md beace v,

+‘m

+(Y+a)PadR

pSRSa

2SVS10

i

Figure 5 Modified Levenberg-Marquardt algorithmEletcher algorithm.

two constants. The rate of convergence was largely insenstive to different choices of p and v,which ledto the ultimate choice of the values p = 0.25 and CT = 0.75. More details of thismodificationisgiven elsewhere [21-261. The optimization procedure updated weights at every connection and yielded rapid androbust training. The weights were initializedto values in the range k0.1 by random assignment. B. StandardError-Back-Propagation Training Routine In a standard back-propagation scheme, updating the weights isdone iteratively. The weights for each connection are initially randomized when the neural network undergoes training. Then the error between the target output and the network predicted output are back-propa-

gatedthrough the network. The back-propagation of error isused as a means to update the connection weights [20]. Repeated iterations of this operation results in a final convergence to a set of connection weights. The general principle behind most commonly used back-propagation learning methodsis the “delta rule,’’ by which an objective function involving squares of the output errors from the network is minimized. The delta rule requires that the sigmoidal function used at each neuronbecontinuouslydifferentiable. This methods identifies an error associated with each neuron for each iteration involving a cause-effect pattern. Therefore, the error for each neuron inthe output layer can be represented as: &,p

=

(Ti,p

-

Si,p)f’(Ui.p)

(8)

Keshavaraj et al.

8

where Ti,pis the desired target output for neuron i and pattern p , and f’, is the derivative of the sigmoidal function used for neuron i. The change in the weight of the connection between neuron i and neuron j is given by:

where P is the learning rate, CY is the momentum factor, and n indexes the iteration. The error signal fromthe neurons in the output layer can be easily identified. This is not so for neurons in the hidden layers. Back-propagation overcomes this difficulty by propagating the error signal backward through the network. Hence, for the hidden layers, the errorsignal is obtained by: 6i.p

= f’(Ui,p)

E (6j.pwiJ) j

(10)

wherej represents the neurons to which neuron i in the hidden layer sends the output. Hence, the weights were updated as shown in Eq. (9).

C. Multiple Adaline (Madaline 111) Training Algorithm Back-propagationandMadaline I11 are useful for trainingmultilayered networks. A gradient-descent technique for minimizing a sum of squared errors is employedinbothof these training routines. Gradientdescent techniques function by finding the lowest error point on any error contour in the weight space. This is accomplished by descending downthe slope or gradient of the contour plane. While back-propagationuses an analytical equation to govern the entire training process. Madaline 111, a network of multiple adalines, uses a more empiricalapproach [27].This approach reduces the system error at each instant in time. Thus, when an input is presented to the network, the resulting output is compared to a target using a sum of the squared error. This approach uses a perturbation method oneach neuron’s summation of inputs to determine the appropriate connection weight. A function called “perturb” is usedin this algorithm. Perturb controls the size of perturbation added to the pre-neuron sumof each neuron.The effect of this added value is propagated through the network. If the network error is reducedas a result of this change, then the perturbation is accepted, and the connection weights are changed. If the error is increased by the perturbation, then a change in the opposite sense is made. The gradient can be used to optimize the weight vector according to the method of steepest-descent:

Here A S is the perturb, p is a parameter which controls

stability and the rate of convergence, and A ( q J 2 / A s is the sample derivative. To some extent, the sample derivative A ( E k ) 2 / A s is Madaline I11 is analogous to the analytical derivative (6&/6sk)2 used in back-propagation. Hence, these two training rules follow a similar instantaneous gradient, and thus perform nearly identical weight updates. However, the back-propagation algorithmrequires fewer operations to calculate gradients than does Madaline 111, since it is able to take advantage of a prior knowledge of the sigmoidal nonlinearities and their derivative functions. Conversely, the Madaline I11 algorithm uses no prior knowledgeabout the characteristics of the sigmoid functions. Instead it acquires instantaneous gradients from perturbation measurements. Finally, any training is incomplete without proper validation of the trained model. Therefore, the trained network should be tested with data that it has not seen during the training. This procedure was followed in this study by first training the network on one data set, and then testing it on a second different data set. VII.

DATA PROCESSING/DATA PREPARATION

It is tempting to view ANNS as simplified versions of biological nervous systems. Yet even the most complex neurocomputers, with several million neurons, are unable to mimic the behavior of a fly, which has approximately one million nerve cells. This is because the nerve system of the fly has far more interconnections than are possiblewith current-day neurocomputers, and their neurons are highlyspecialized to perform necessary tasks. The human brain, with about 10 billion nerve cells, is stillseveral orders of magnitude more complex. Even so, artificialneural networks exhibit many “brainlike” characteristics. For example, during training, neuralnetworks may construct an internal mappingl model of an external system. Thus, they are assumed to “make sense” of the problems that they are presented. As with any construction of a robust internal model, the external system presented to the network must contain meaningful information. In general the following anthropomorphicperspectives can be maintained while preparing the data: Determine whichdata are tobe used to train and test the selected network. For effective training, the selected data should be relevant, meaningful, and complete. Divide the available data into training and test data sets (1/3). Test sets are used to validate the trained network and insure accurate generalization. Scalelmap the data to an array of numbers. Transform the real world data into numeric input and target output patterns.


Normalize both the input andthe target output data to fit the transfer function range. This implies that the data have to be scaled to fit betweenthe minimum and maximumvalues of the selected transfer function.

of Experimental/ReaI-World

A.Scaling Data

In many cases, numeric data with a continuous range of values do not need to be encoded prior to being sent to a neural network. In the authors' experience encoding was found to be necessary only in some special cases. In such cases, the analog data had to be scaled to fit within the transfer function range. A maximum-minimum scaling process for standardizing numericdata was used. To do this, the following formula was applied to each data point:

1

X (TFMAX- TFMIN)+ TFMIN

where X, is the Max/Min scaled data point; x d is the experimental unscaledinput data point; &AX is the minimum value of the raw input data; d M I N is the maximum value of the raw inputdata; TFM, is the transfer funcTFMINis the transfer function tionmaximumand minimum. Similarly, the network predicted data must be unscaled for error estimation withthe experimental output data. The unscaling was performed using a simple linear transformation to each data point.

X

M MAX -

1

~MIN)

+

~MIN

where Y d is the actual experimental outcome; YS is the neural network predictedoutput; 6" and ~ M I Nare the maximum and minimum value ofthe raw output data. VIII. OPTIMAL TRAINING

During the training phase, one of the important objective is to achieve optimal training. When training is optimal, the number of training examples required to assimilate the cause-effect relationship between the variables modeledisminimized. The goodness of the internal model built by the neural network depends on this efficiency.Severalmethods to achieve optimaltraining have been developed by earlier investigations [15,16,28]. A novel approach, suggested by Weigandet al. [29],uses

9

a separate validation data set during training. A validation data set is made up of arbitrary number of data points from the original training data set (for example, 10%). Hence, at the end of each exposure, after updating the weights, the validation data are presented to the network and the network's prediction error for the validation is calculated. Training is stopped when this error starts to increase. However, it should be remembered that validation data are not part of the training data once the training begins. IX. HOW MUCH DATA, NEURONS, AND LAYERS ARE NEEDED WITH ANNs?

In neural network design, the above parameters have no precise number/answers because it is dependent on the particular application. However, the question is worth addressing. In general, the more patterns and the fewer hidden neurons to be used, the better the network. It should be realizedthat there is a subtle relationship between the number of patterns and the number of hidden layer neurons. Havingtoo few patterns or too many hidden neurons can cause the network to memorize. When memorization occurs, the network would perform well during training, but tests poorly with a new data set. An easy to apply rule of thumb that the authors have found useful intheir experience to determine the number of patterns required for training was: Training patterns

+ hiddens + outputs) to no. of (inputs + hiddens + outputs) (14)

= 2 X no. of (inputs

10 X

where the inputs, hiddens, and outputs indicated the number of neuronsin those layers of a selected architecture. This equation suggests that the number of training data required will be between 2 to 10 times the number of neurons in the selected network. It should be emphasized that the selected network need not learn every training pattern perfectly. A good network will generalize wellfor new testing cases. Many users make the mistake of trying to achieve perfect training by adding more and more hidden neurons without realizing that this only leads to a higher possibility of deterioration during testing. There are no hard and fast rules to overcome this problem,but either of the following two guidelines can be used. Number of hidden neurons

- (inputs + outputs) or 2 Minimum number

of hidden neurons

(15)

- [number of patterns] - inputs - outputs 10

Keshavaraj et al.

10

Maximum number of hiddenneurons [number of patterns] - inputs - outputs 2 During the selection of the number of hidden layer neurons, the desired tolerance should also be considered. In general, a tight tolerance requires that the selected network be trained with fewer hidden neurons. As mentioned earlier, cross-validation during training can be used to monitor the errorprogression, which subsequently serves as a guideline in the selection of the hidden layer neurons. A network that is too large mayrequire a large number of training patterns in order to avoid memorization and training time, while one that is too small may not train to an acceptable tolerance. Cybenko[30] has shown that one hidden layer with homogenous sigmoidal output functions is sufficient to form an arbitrary close approximation to any decisions boundaries for the outputs. Theyare also shown to be sufficientfor any continuous nonlinear mappings. Inpractice, one hidden layer was found to be sufficient to solve most problems for the cases considered in this chapter. If discontinuitiesin the approximated functions are encountered, then more than one hidden layer is necessary.

X. APPLICATION OF ANNs IN CHEMICAL SCIENCE Applications of neural networks are becoming more diverse in chemistry [31-401. Some typical applications include predicting chemical reactivity, acid strength in oxides, protein structure determination, quantitative structure property relationship (QSPR), fluid property relationships, classification of molecular spectra, group contribution, spectroscopy analysis, etc. The results reported in these areas are very encouraging andare demonstrative of the wide spectrum of applications and interest in this area.

relationships have been well addressed. In addition to this rather simple case, a relatively new application of neural networks namely in QSPR is described in Case V. The overall results obtained from these examples havebeenveryencouragingand demonstrate that ANNS should beconsidered as a potential tool in polymer science. A. Case I: Fluid Property Prediction of Siloxanes An artificial neuralnetwork based approach for modeling physical properties of nine different siloxanes as a function of temperature and molecular configuration will be presented. Specifically, the specific volumes andthe viscosities of nine siloxaneswere investigated. The predictions of the proposed model agreed well with the experimental data [41]. 1. Materials The compounds considered in this work are shown in Table 1. These particular siloxanecompoundshave been described elsewhere [42]. A system of abbreviation already in use for these compounds is given in the same table. Their abbreviated formulas are expressed on a standard functional basis inTable 1 along withthe name of the compound. An explanation of the units in the abbreviated formulas is given below: 0 0

M, the monofunctionalunit:(CH3)3Si-0

0 0

D, the difunctional unit:

0-

0 0

(16)

CH3 1

:

Si 4

'I:

(17)

0

CH3

XI. CASE STUDIES In all, five case studies are presented in this chapter in an attempt to illustrate the performance of applicability of neural networks in polymerscience. The diversity of problems considered in this chapter highlight ANNS as a general tool for problem solving. Although there are numerous other methods that may be just as appropriate for problem solving, ANNS are rapidly becoming the method of choice in certain areas. This is particularly so in those areas where the phenomenological understanding is very limited [23]. This section begins witha very simplecase; that of modeling the specific volumes and viscosites of siloxanes. The siloxanes, which are considered here, have been well documented the in literature and the structural

2. Specific Volumes of Siloxanes Specific volumesof polymers generally show a regular decrease with increase in temperature, for both chain and ring structures. Experimental data from earlier investigations [42] indicate that this decrease in specific volume is greater for chain compounds than for ring compounds. For each increase of one difunctional group in the molecule, the specific volume of the chain compound is usually greater than that of a ring compound containing the same number of silicon atoms. It is hypothesized that the ring structure is more compact, and contains stronger interatomic attractive forces holding the molecule than does the chain compounds. With increased molecular size, this type of differentiation becomes much more complicated.

11

Artificial Neural Networks as a Modeling Tool Table 1 List of Siloxanes and Their Abbreviations

Name Abbreviation

No.

of siloxane

MM MDM MD2M MD3M MD4M

1. 2. 3. 4.

5. 6. 7. 8. 9.

Hexamethyldisiloxane Octamethyltrisiloxane Decamethyltetrasiloxane Dodecamethylpentasiloxane Tetradecamethylhexasiloxane Octamethylcyclotetrasiloxane Decamethylcycloentasiloxane Dodecamethylcyclohexasiloxane Tetradecamethylcycloheptasiloxane

D4

D5 Ds D7

The specific volumes of all the nine siloxanes were predicted as a function of temperature and the number of monofunctional units, M, and difunctional units, D. A simple 3-4-1 neural network architecture with just one hidden layer was used. The three input nodes were for the number of M groups, the number of D groups, and the temperature. The hidden layer had four neurons. The predicted variable wasthe specific volumes ofthe silox-

anes. A schematicof the typicaltraining process is shown in Fig. 6. The experimental specific volume data were available in the temperature range of 273K to 353K, with 20K increments. The nine types of siloxanes were arbitrarily dividedinto two groups,one each for training and testing. The compounds 1, 2, 4, 6 , and 8 were utilized in the trainingphase. The trainednetworkwasthen

Number of monofunctional units (M) Number of difunctional units (D) Temperature

l-

Neural network model

n

I

r

Neural network model predicted specific volume

Adjusted weights

Experimental

Nonlinear optimization routine

specific volume of siloxanes

Figure 6 Schematic of a typical neural network training process. I-input layer; H-hidden layer; 0-output layer; B-bias

neuron.

Keshavaraj et al.

12

tested for the specific volume predictionof compounds 3,5,7, and 9. The training andtesting results are shown in Fig. 7. The root mean squared (rms) error for the model was 0.003.

where v is the specific volume of the siloxane, is the viscosity of the siloxane and, C and W are constants that can be obtained from experimentaldata. Values of these constants for the siloxanes under investigationwere reported by Hurd [42].It should be remembered that C and W are a function onlyof the compound and independent of temperature. The experimentalviscosity data for the siloxanes under investigation were available at 298 K, 323 K, and 348 K. The proposed neural network model was put to a test by comparing its predictions with Batschinski’s equation in this temperature range (298K, 323 K, and 348 K). Figure 9 shows the comparison betweenthese two predictions. The agreement was very good between the experimental based Batschinski’s relationship and the ANN model at least within the limited range of training data.

3. Viscosities of the Siloxanes Viscosities of the siloxanes were predicted over a temperature range of 298-348 K. The semi-log plot of viscosity as a function of temperature was linear for the ring compounds. However, for the chain compounds, the viscosity increased rapidly with an increase in the chain lengthof the molecule. A simple 2-4-1neural network architecture was usedfor the viscosity predictions. The molecular configuration was not considered here because of the direct positive effect of addition of both M and D groups on viscosity. The two input variables, therefore, were the siloxane type and the temperature level. Only one hidden layer with four nodes was used. The predicted variable was the viscosity of the siloxane. For all the siloxanes the network wastrained at two temperature levels; 25°C and 75°C. The trained network was then tested for its viscosity predictions at 50°C.The network training andtesting results are shown in Fig. 8. The rms error for this prediction was 0.002.

B. Case II: Prediction of Densities of High Molecular Weight Esters Usedas Plasticizers*

To any one concerned with the production of plasticized resins, compatibility has long meant the ability of two or more materialsto mix with eachother to form a homogeneous compositionof useful desired plastic properties [44,45].A plasticizer, therefore, can be definedas a material that will soften and make inherently rigid andeven brittle polymers flexible. Plasticizers are also sometimes

4 . RelationshipBetweenViscosityand Specific Molar Volumes Batschinski’s relationship [43]for specific volume and viscosity is probably the best known for siloxanes. This relationship can be written as follows: v =

(g) +

* Reprinted

W

with permissionfrom Adv. Polyrn. Tech., Vol. 14, No.

3, 215-225 (1995). 0 1995 by John Wiley BC Sons, Inc.

Neural network mining Experimental data 0

0.7

Neural network testing

0.8

0.9

I

1.1

I .2

1.3

Experimental specific volume(ml)

Figure 7 ANN training and testing results for specific volumes of siloxanes.

I .I


Neural network

m

13

mining

E.\perimentaldata Neural network testing

0

2

0

.

8

6

1

12

10

Experimental viscosity (cp) Figure 8 ANN training and testing results for viscosities of siloxanes.

3

1

0.

1.4

Batschinski Equation

.1 1.3 n E V

Specific volume at 323K 0

g

Specificvolumeat298K

J

Specific volume at 348K

-

0.9 0.8

I

0.8

0.9

1

1.1

1.2

1.3

1.J

Specific volume predicted by Batschinski equation (ml) Figure 9 Comparison of viscosity predictions betweenANN model and Batschinski's equation.

1.5

Keshavaraj et al.

14

added to lower the melt viscosity. They can reduce the temperature of the second-order transition for easier processing by virtue of both glass transition temperature ( Tg)depression and throughinternal lubrication. Hence, plasticizers are commonly liquids that are compatible with their parent polymers, but notto the point of complete miscibility. Primary plasticizers that are used as thesole plasticizing ingredients are typically high molecular weight organic esters. Such esters have found extensive use in the plastic industry. They demonstrate a high degree of permanence and flexibility ona unit weight basis. Also of importance are their low levelof toxicity, consistently light color, resistance to degradation, and relatively low a reduction in effectivecost. Plasticizers generally show ness with increasing molecular weight and temperature. As solvents they give solutions that undergo gelation on heating. Lower molecular weight esters, however, can become volatile and/or decompose at elevated processing temperatures. Since the physical properties of the polymer matrix can be significantly modified by the type of the plasticizer used, these high molecular weightesters have been extensively investigated during the last decade. Liquid density, for example, has been foundto be a useful bulk property of these liquid plasticizers, since density helps define the internalmolecularflexibility.Most of the time, however, experimental difficulties prevent the accurate determination of density. Moreover, existing empirical modelsto predict density remaintedious because of many ill-definedparameters. Therefore, a simple neural network based model approach is offered wherein the effect of the number of end groups andtemperature on density were incorporated. The proposed model is also capable of differentiating noiseor error in the experimental data. The effects of structure and temperature on the densities of a homologous series of esters were investigated. In all, five different series of esters were considered.

They were: dialkyl sebacate ester (DSEs), 1,IO-decanedio1 diester (DDEs), triglycerides (TGEs), 2-(hydroxymethyl)-2-methyl-l,3-propanedioltrialky esters ( R E S ) , and pentaerythritoltetra alkyl esters (PTEs). Of the five, the first two can be characterized as linear and flexible esters, while the thirdand fourth are “Y” branched compact molecules. The last series, namely PTEs are “ + ” shaped and highly rigid. The various esters considered in this study were structurally differentiatedby earlier investigators[46-481 who usedthe number of methylene groups in the molecule.Density measurements should help reveal changes in the compatibility and efficiency of the ester, and relate directly to the plasticizer’s permanence. The results of this study offer an insight into the density changes that are produced by structure for these industrially important plasticizers. 1. Materials The isothermal experimentaldensity data for all the five type of esters were obtained from various literature sources [46-481. The reported density measurements for DDEs, TGEs, and PTEs were for different temperatures (310-413 K). The chemical-structures of all five esters are shown in Tables 2-6, with the different numbers of methylene groups in the molecules being specified by (X).Numerical results from these references are not presented here. The isothermal densities within each of the five groups of esters reveal a significant dependence on the molecular structure. At a constant temperature and pressure, the liquid density of the ester decreases with increasing molecular weight and the successive addition of the methylene groups, ( X ) , even though the concentration of the end-groups remained constant. It should be remembered that the volume requirement for the rotation of a molecule depends on both its shape and the possible directional nature of its interaction with its neighbors. As the chain length inthe ester is increased by the successive addition of methylene

Table 2 Dialkyl Sebacate Esters (DSEs) 0

I

0

I

CH3(CHz)+C+CHz)rC+CHz)xCH3

Compound weight

Molecular

314 Di-n-butyl sebacate 370 Di-n-hexyl sebacate 426 Di-n-octyl sebacate Di-n-decyl sebacate C34H6604 Di-n-dodecyl sebacate C38H7404 Di-n-tetradecyl sebacate C42H8204 Di-n-hexadecyl sebacate C46H9004 Di-n-octadecyl sebacate

X formula 3 5 7 9 11 13 15 17

Empirical C18H3404 C22H4204 C26H5004 C30H5804

482

65 1

Neural

Artificial

a

Networks as

Modeling Tool

15

Table 3 1,lO-DecanediolDiesters(DDEs)

Compound weight

Molecular

X formula 2

1,lO-decanediol dibutyrate 314 1,lO-decanediol dioctanaote 426 1,lO-decanediol didecanoate 482 C42H8204 1,lO-decanediol dipalmitate

6 8 14

groups, it tends to loosen the packing by disrupting the large sterically hinderedcarboxyl group in the ester series. Thisdisruption produces an increase in the specific volume of the ester, hence, the packing efficienciesfor small molecules are higher than they are for the larger molecules. The densities of individual groups of esters can also be differentiatedby the differencein the number of carboxyl groups. The observed decrease in the density of the esters, with the successive addition of the methylene groups(X),was probably due to the dilution of the X O O - groups. This is because the density of the carboxyl group isgreater than the methylene group. The densities of the esters were in the following order: PTEs > TTEsEGEs > DSEs > DDEs

This observed order can be explained on the basis of

Empirical C18H3404 C26H5004 C30H5804

the relative degree of the packing and flexibility of the atoms in the molecules of these esters.

2. Neural NetworkPerformance A very simple 2-4-1 neural network architecture with two input nodes, one hidden layer with four nodes, and one output node was used in each case. The two input variables werethe number of methylene groups and the temperature. Although neural networks have the ability to learn allthe differences, differentials, andother calculated inputs directly fromthe raw data, the training time for the network can be reduced considerably if these values are provided as inputs. The predicted variable was the density of the ester. The neural network model was trained for discrete numbers of methylene groups over the entire temperature range of 300-500 K. The

Table 4 Triglyceride Esters (TGEs)

Compound weight

Molecular X

Tributyrin Tricaproin Trioctanoin Tridecanoin Trilaurin Trirnyristin Tripalmitin Tristearin

2 4

6 8 10 12 14 16

formula

Empirical 302 386 470 554 639 723 807 890

Keshavaraj

10

Table 5 2-(Hydroxymethyl)-2-Methyl-l,3-Propanediol Triesters (TTEs) 0

II

CHZ”C”tC&)&H3 0 H3C-

I

II

C--CH*+CH~)&HJ 0

II

CH*+CHz)&H3

Compound

X

2-(Hydroxymethyl)-2-methyl-1,3-propanediol tripropionate 2-(Hydroxymethyl)-2-methyl-1,3-propanediol tributyrate 2-(Hydroxymethyl)-2-methyl-1,3-propanediol trivalerate 2-(Hydroxymethy1)-2-methyl-l,fpropanediol triheptanoate 2-(Hydroxymethyl)-2-methyl-1,3-propanediol tripoctanoate 2-(Hydroxymethyl)-2-methyl-1,3-propanediol trinonanoate 2-(Hydroxymethyl)-2-methyl-1,3-propanediol tridecanoate 2-(Hydroxymethyl)-2-methyl-1,3-propanediol tripalmitate

Empirical formula

Molecular weight

1

C14H2406

288

2

5

C17H3006 C20H3606 C26H4806

330 372 456

6

C29H5406

498

7

C32HaoO6

540

8

&H6606

582

C53Hl02O6

834

3

14

Table 6 Pentaerythritol Tetra Alkyl Esters (PTEs) 0

II

CHZ”C+CHz)dHs

l

0

II

0

II

C H ~ C H ~ ) ~ U H Z C ~ ~ H ~ - ~ - ( C H ~ ) ~ H ~ 0

II

CHz-”-C--(CHz)dH3

Compound weight Pentaerythritol Pentaerythritol Pentaerythritol Pentaerythritol Pentaerythritol Pentaerythritol Pentaerythritol Pentaerythritol

Molecular

tetraacetate tetrapropionate tetrabutyrate tetravalerate tetraheptanoate tetraoctanoate tetranonanoate tetradecanoate

Xformula 0 1

2 3 5 6 7 8

Empirical CnH200s CnH2sOs c21H3608 CzsH.1408 C33HaoOs C37H6808 C4lH7608 C45HS408

304 370 416 472 584 640

696 752

““x” at training

“X” at testing

~u~-of-s~uared-e~or

3, 7, 11, 15 and 17

2, 6, 10, 14 and 16 1, 3, 6, 8 and 16 1, 2, 4, 6 and 8

trained network was then tested for density predictions using compounds with a d i ~ e r e nnumber t of methylene groups [49]. The latter were not included in t ode1 are given in Table 7 , pounds used in the training and testing phases. Each ester group was trained an tely. The s u ~ - o f - s ~ u a r eerror d of the prowork for the density predictions during the was very low (Table 7). The density predicith a high degree of recision consi~ering

4, 8 and 12 2, 5 and 7 3 , 5 and 7

Na > K, and their hydroxides are arranged in the sequence KOH > NaOH > LiOH as to their inhibitive efficiency relative to thermal destruction of polyolefins. And the efficiency of alkali metals can be represented by the sequence Na > K > Li. This seems to be due

4

= 35,000.

to the fact that the electron-transfer work of sodium is somewhat higher than that of potassium, whereas the atomic dimensions of sodium are much less than those of potassium, and sodium radical is more effective accepting the macroradicals formed upon thermal destruction of polyethylene as compared to the potassium radical. The inhibitive efficiencyof alkali metalhydroxides increases withincreasedbranching of polyethylene. This is confirmed by more pronounced effect of these hydroxides diminishingthe yield of propane and propylene than in case of ethane and ethylene. The decreased yield of propane and propylene is also conditioned by more efficient inhibitionof the macroradical isomerization stageby alkali metalhydroxides. Upon thermal destruction of polyethylene with the use of inhibitors the

Stabilization High-Temperature

87

numberofvinylidene (882 cm") andvinylene (965 cm") bonds is drastically growing, while the number of vinyl bonds (909 and 993 cm") decreases sharply. The overall nonsaturation of destructed polyethylene with the use of an inhibitor isincreased by a factor of 1.5

A-

-CHdH

as compared to the initial one, but branching is greatly reduced. At the first stage of polyethylene thermal destruction the metallizing of polyethylene macroradicalby the metal radical takes place. H

+

Me.+ - C H d H d - M e

AH3 One would think that thermal destruction of polyethylene should be inhibited byhydroxides of alkali metals according to the following scheme, as with phenols: H

+

- C H d H d -

AH3

H

MeOH+ - C H d H d H

AH3

+

MeO.

AH3

H

I

+

-CHdH&.

MeO.

AH3 However, this mechanism is at variance with the following factors: (1) high bonding strength of OH-Me (90 kcalhol) that is much greater than that of OH in phenolmolecules thus eliminating the possibility of breaking at temperatures from 350 to 400°C; (2) the absence of the absorption bonds 700-900 cm" characteristic of the bonding C U M e in infrared spectra of

1

C H d H

- C H 4 H 4 H Z

polyethylene after destruction; and (3) the absence of EPR-signal in the products after destruction. Metallizing issupported by the fact that thermal destruction of polyethylene is inhibited by alkali metals. At the second stage the decomposition of metal alkyl takes place. Metal alkyl is liable to decompose into metal alkyl and olefin causing the increased saturation of uolyethylene macromolecules:

-

AH3

+-

CHdHdHT"e

+

+ CHHH"CHdH2-

Metal alkyl is liableto decompose with breakingof the carbon-carbon bond that is in P-position relative to the C-Me bond: -CHdHdH-CH&H3

I Me

+ R-Me

+

CH4H-CHdH2\

where

The resultant metal alkyl is rapidly decomposedat temperatures from 200-250°C with the formation of This reaction leads to drastic decrease in the number metal, alkanes, and alkenes C 4 3 . of methylgroupsin polyethylene and to an increase inMetalalkylisliable to react with nonsaturated ends the number of double bonds. of polyethylene molecules: R=Cd,.

-CHdHdH=CH2

+

r

n -CHdHz-+ AH3

J

C H d H -CHdH-CH2HI

AH3

Losev

88

As this takes place, scission of polymeric chains is hindered andthe molecular dimensionsare even growing up to the temperature of 300°C.

"!

-C H d H d H -

+

€J

M

Metal alkylreacts with nonsaturated ends of macromolecules followingthe pattern of substitution reaction [and causinga decrease in the number of methyl groups]:

-CHdH2-+

CHz-C=CH

'A

AH3 111.

H

INHIBITION OF THERMAL DESTRUCTION BY BORIC ACID ESTERS

Boric acidesters have aroused considerable interest because they are stable up to the temperature of 450°C, display low volatility and good compatibility without migration from olefin. Boric acid esters provide for thermal stabilization of low-pressure polyethylene to a variable degree (Table 7). The difference in efficiency derives from the nature of polyester. Boric acid esters of aliphaticdiols andtriols are less efficient thanthe aromatic ones. Among polyesters of aromatic diols andtriols, polyesters of boric acid and pyrocatechol exhibit the highest efficiency. Boric acid polyesters provide inhibitionof polyethylene thermal destruction following the radical-chain mechanism, are unsuitable for inhibition of polystyrene depolymerization following the molecular pattern and have little effect as inhibitors of polypropylene thermaldestruction following the hydrogen-transfer mechanism. The inhibitive efficiencyof boric acidpolyesters differs greatly. The highest efficiency is exhibited by polyesters of boric acid, aromatic diols and triols. This derives from the fact that in this case the radicals are accepted not only by boron, but also by the aromatic nucleus. Among the aromatic polyesters, most efficient is ester of boricacidand pyrocatechin due to the Frank-Rabinovich cage effect. The efficiency of inhibi-

+

CH3-CHAH2-

H2Me

tion in case of polyethyleneis cymbatically growing with increase in branching, being associated with growing importance of the disproportionationand macroradical-decomposition reactions as compared to the chain transfer reactions, since branchingpresents hindrance to the detachment and transfer of hydrogen to the macroradical. In presence of polypyrocatechin borate the branching of polyethylene after destruction is greatly extended, whereas nonsaturation is growing only slightly.Interaction between polypyrocatechin borate and macroradicals of polyethylene gives inhibiting radicals that have been detected by the EPR method. After heating under the same conditions, but without an inhibitor the samples of polyethylene revealed no EPR-signal. In the process of inhibition polypyrocatechinborate interacts with polyethylene macroradicals to form the B " C bonds. This is confirmed by the fact that the absorption spectrum of polyethylene inhibited with polypyrocatechin borate revealed the bands in the region of1350cm" characteristic for the B " C bond. There is nosuch a band inthe spectrum of pure polypyrocatechin borate after heating under the same conditions. Chemical analysis of boron in polyethylene provides support for the IR-spectroscopy data concerning the presence of chemically bonded boron in polyethylene after destruction. Inhibition of polyethylene thermal destruction by polypyrocatechin borate could be represented as follows. The initial molecular-chain scission of branched

Table 7 Inhibition of Polyethylene Thermal Destruction by Boric Acid Polyesters

Polyester of Ethylene glycol Glycerol Pyrogallol Hydroquinone Resorcinol Pyrocatechol Pyrocatechol Polypropylene

Molecular weight after destruction

Molecular weight before inhibition Without destruction

inhibitionWith

1050 1050 1050 l050 1050 2650 2100 6300

158,000 158,000 158,000 158,000 158,000 34,800 3 1,500 366,000

Temperature = 420°C; Time = 2 hours; Stabilizer concentration

=

5%.

9200 15,100 20,500 16,800 14,200 24,000 15,000 22,200

High-Temperature Stabilization of Polyolefins

89

polyethylene leads to the formation of radicals RI and R2 : H - C H d H d H d I< H d H d H

+

CH3

+

- C H d H d H d H -

AH3 Isomerization, disproportionation, and decomposition reactions of the radical R1 yield nonsaturated end groups C H d H " C H 2 and result in the formation of ethane and ethylene. Isomerization and decomposition H R2 ( - C H d H d * )

-+

+

CHdH=CH2

AH3 H I

CHdHdHd"CH2(!!H3

of the radical R2 proceed withthe formation of propylene and new end radicalthat reacts with the macromolecule of polyethylene according to the following scheme:

-CHdH

I

H

As a result, the central radical R3 is formed, and the fragment with the end methyl group breaks down into propane and a new fragment with the end vinyl group. Polypyrocatechin borate has a higher inhibitiveefficiency relative to the reactions of the radical R2 than to that of the radical R,, since the latter exerts a more severe decrease in the yield ofpropane and propylene than in the yield of ethane and ethylene, and favors an insignificant growthof polyethylene nonsaturation in the pro-

As this takes place, the inhibitor molecule is regenerated, and side methyl group is substituted by end vinyl [411.

IV. STABILIZATION OF CURED POLYETHYLENE Polyethylene cured by the chemical and radiation-chemistry methods undergoes thermal destruction upon heatingas in normal polyethylene. Thermostabiliz-

cess of destruction. However, this offers noexplanation for the diminished number of methyl groups in polyethylene as compared to polyethylenewithout the use of inhibition. It is believedthat the recombination product is liable (due to the increased mobility ofthe hydrogen atom following the carbon atom in P-position relative to the C-"C bond) to isomerizationanddecomposition along the ester bond:

ers should possess good thermal stability up to 300-350°C being no bar to chemical vulcanization, i.e., being inert to the curing agent. Besides, these thermostabilizers should have good compatibility, low volatility without sweating out from the polyethylene mass. Currentlyused thermostabilizers of polyethylene (alkyl phenols, aromatic amines and diamines, sulphur- and phosphorus-containing stabilizers) exhibit thermal stability up to the temperatures from 200-250°C, decompose at higher temperatures and contribute to the initiation of thermal destruction. They react with a curing

90

Losev

fraction becomes 24% at 350”C, 41% at 380°C, and 92% at 400°C. [47]. An investigation into the effect of the concentration of polyaniline disulphide on inhibition of thermal destruction in case of cured polyethylene has demonstrated that polyaniline disulphide is efficient even at the concentration of 0.25%. An increase in the concentra[Q]. At present, very few compounds are used as thertion over the range 0.25-1.0% results in the increased mostabilizers for cured polyethylene. Among them may efficiency, while further increase in the concentration be listed 1,3-dihydro-2,2,4-trirnethylquinoline;P-dileads to a slight drop in inhibition. naphtyl, n-phenylenediamine, zinc mercaptobenzimidaPolyamine disulphides as inhibitors of thermal dezole [43-451. struction of cured polyethylene are effective over a long Polyamine disulphides do not inhibit peroxide vulperiod of time. canization of polyethylene, are stable in air up to Upon 10-h exposure, a decrease ingel fraction 300-350°C, exhibitgoodcompatibilityandshowno amounts to .40% without the use of inhibition andto 8% sweating out from the polyethylene mass. Table 8 gives in presence of polyaniline disulphide. High efficiency of the comparison betweenthe efficiency of polyamine dipolyaniline disulphide has been confirmed by the investisulphides as thermostabilizers of cured polyethylene. gation of strength properties for cured polyethylene. A series of polyamine disulphides (polyaniline disul- Tensile strength of cured polyethylene in vacuum at phide, polyamine disulphide, and polyparaphenylenedi300°C decreases insignificantly, and after a IO-h expoamine disulphide) represent effective thermostabilizers sure it amountsto 192 kg/cm2. Without inhibition tensile of cured polyethylene, and provide a decrease in gel strength decreases depending onthe exposure time, and fraction 2.5-3 times as large as that in case of inhibited over a period of 10 hours it comes to only 58 kg/cm2. thermal destruction. Stabilizers of normal polyethylene With the use of polyaniline disulphidea specific elonga(Neozone “D”, Santonox “R”) are inefficient as stabition at rapture decreases slowly as a function of time, lizers of cured polyethylene, these substances decomand after 10 hour exposure this parameter amounts to pose andeven initiate thermal destruction of cured poly90% of the initial value. Withoutan inhibitor a specific ethylene. elongation decreases sharplywith increase in time, Polyamine disulphidesare effective thermostabilizamounting to only 40% of the initial value after 10-h exposure. Polyaniline disulphide isan efficient inhibitor ers of cured polyethylene up to 400°C. In presence of of heat aging in case of cured polyethylene (Table 9). polyamine disulphides a decrease in gel fraction is one half as large as that of nonstabilized cured polyethylene A drop in gel fraction of nonstabilized cured polyover the temperature range from 350-380°C [46]. ethylene amounts to 50% after 25-h exposure, 75% after 50 hours and after a 75-h exposure the complete fall is The highest efficiency is exhibited by polyaniline observed. At the same time, a decrease in gel fraction disulphide in presence of which a decrease in gelfraction in presence of polyaniline disulphide is observed only amounts to 7% at 350”C, 19% at 380”C, and 60% at 400”C, whereas without an inhibitor the decrease in gel after a 50-h exposure and comes to only 2%, whereas

agent (most commonly peroxide) causing reciprocal consumption of peroxide and thermostabilizer. To take an example, the use of P-dinaphtyl n-phenylene diamine (diafen N N ) as a stabilizer in compositionswith dicumyl peroxide involves a decrease in gel fraction by 15-40% depending on the concentration of dicumyl peroxide

Table 8 Inhibition of Thermal Destructionof Low-Density Cured Polyethylene in Vacuum

(10-3 torr) Decrease in gel

Gel fraction (%) Before destruction destruction After fraction Polyaniline disulphide Polyparaphenyldiamine disulphide Polyamine disulphide Polycarbamide disulphide Polythiocarbamide disulphide Polyethylenediamine disulphide Without stabilizer Neozone “D” Santonox “R” Temperature

=

84

82

64 70 59 70 87 78 75

380°C; Inhibitor concentration = 1%; Time = 120 min.

68 60 50 44 47 52 43 49 39

19 27 22 39 24 26 41 62 49

91

High-Temperature Stabilization of Polyolefins Table 9 Heat Aging of Cured Polyethylene in Air ~

Decrease in gel fraction (%) Time (h) inhibitor With inhibitor Without 25 50 75 100

76 100

-

0 2

5 10

Temperature = 150°C; Polyanilinedisulphideconcentration 0.8% by mass.

=

after 100 hours a decrease in gel fraction amounts to 10% [48]. High thermostabilizing efficiency of polyamine disulphides relative to chemically cross-linked polyethylene is conditioned by the abilityaccept to macroradicals at the disulphide bridge and imine group. Besides, the presence of paramagnetic centers causes the adherence of macroradicals providingfor an extra stabilizingeffect

W].

REFERENCES 1. A. G. Sirota, ModiJicationof Structure and Properties of Polyolefins, Chemistry Publishers, Leningrad, p. 3 (1994). 2. A. K. Wardenburg, Plastics in Electrotechnical Industry, Gosenergoizdat Publishers(1957). 3. V. S. Shifrina and N. N. Samosatsky, High-pressure Polyethylene, Goschimizdat Publishers(1958). 4. S. S. Mindlin and N. N. Samosatsky, Production ofPoly-

5. 6. 7. 8. 9. 10. 11.

ethylene Products by the Extrusion Method, Goschimizdat Publishers (1959). Low-Pressure Polyethylene (N. M. Egorov, ed.), Goschimizdat Publishers (1958). F. Klema, Mitt. Chem. Forsch., 12: 159(1958). H. Stepanek, Plastverarbeiter, 10, 137(1959). Madorsky, Straus, Thompson, and Williamson, J . Polymer. Sci., 4: 639 (1949). Hopff Kunststoffe 42: 2;423-426(1952). H. U. Ellinek, J . Polymer Sci., 4: 850(1948). R. Simka, L. A. Wall, and R. G. Blatz, J . Polymer Sci.,

5: 615 (1959). 12. L. A. Wall and S. Straus, J . Polymer Sci., 44: 113 (1960). 13. N.N. Semenov, On Some Problemsof Chemical Kinetics and Reactivity, USSR Ac. Sci. Publishers (1959). 14. I. Feugt, Stabilization of SyntheticPolymersAgainst

Light and Heat Effect, Chemistry Publishers, Leningrad, p. 26-29 (1972). 15. I. Feugt, Stabilization of SyntheticPolymersAgainst Light and Heat Effect, Chemistry Publishers, Leningrad, p. 184-185(1972). 16. A. G.FarbenfabrikenBayer, Neth. Appl., 6: 5, 15,965 ( 1966). 17. French patent 1350966,23, XI1 (1969). 18. Badische Anilin Soda-Fabrik A . G. To Adolf-Hrubeschanol, Hans Moeller Belg., 617,190, 5 , XI (1962).

19. E. Glazenceand F. Holstrup,EastmannKodak Co., Brit., 972989, 21, X (1964). 20. Eastman Kodak Co., (C. E.Fholstrup, ed.), France I, 366, 533, 10 VI1 (1964). 21. R. Renneth, Hills and R. G. Walton, Phillips Petroleum Co., U.S. 3, 227,676 (1969). 22. Azien de Colori National AfJiniof CNASPA (G. Vigailo, ed.), Ital., 630562, 18 (1961). 23. Badische Anilin Soda-Fabrik A . G. (HansBurgerand BeinhurdRazinalla, eds.), German I, 159,645,19 XI1 (1965). 24. Badische Anilin Soda-Fabrik A.G. Neth. Appl. 6, 413, 754, 31 V (1969). 25. MontecatiniSocieta General1 per Industria Mineraria Chemica Ital. 642, 042, 5 VI1 (1962). 26. Hitto Chemical Industry Co. Ltd.(F. Scki, J. Kanakami, and F. Katsamura, eds.), Japan 645, 5166, 19 IV (1966). 27. British patent 833853(1960). 28. Shell Internationale Research (T. H. Bouthle and C. C. Gosselik, eds.), German, I, 203, 583, 9 XI1 (1965). 29. A. V. Topchiev, Ya. M. Paushkin, A. V. Nepryakhina, P. G. Anan’ev, andN.N. Dmitrievsky,Reports of USSR Ac. Sci., 133: 134 (1960). 30. A.V. Topchiev, Ya. M. Paushkin, A.V. Nepryakhina,

P. G. Anan’ev, and N. N.Dmitrievsky, Zzvestiya USSR Ac. Sci., Chemical Series, 10: 1838(1960). 31. Ya. M. Paushkin, Yu. P. Losev, andP.G. Anan’ev, J . Nefekhimiya, I X 60-62 (1969). 32. Ya. M. Paushkin, Yu. P. Losev, and P. G. Anan’ev, IzvestiyaUSSR Ac.Sci., ChemicalSeries, 6: 1276-1278

(1969). 33. Yu.P. Losev, Ya. M., and Paushkin, V.M. Khoruzhy Reports USSR Ac. Sci.,XVIII, 1014-1015 (1974). 34. Yu. P. Losev, Ya.M. Paushkin, V. M. Khoruzhy, and D. I. Metelitsa, ReportsUSSR Ac. Sci., 218: 390-392 (1974).

35. Yu. P. Losev, Ya. M. Paushkin, V. M. Khoruzhy, and G . V. Dedovich, Reports USSR Ac. Sci.,218: 1365-1367

(1974). 36. Yu. P. Losev, and Ya.M. Paushkin, Reports Belarussian Ac. Sci. XII, 522-525 (1968). 37. Ya. M. Paushkin,and Yu.P. Losev, J . Polymer Sci., Part C , 501-511 (1968). 38. Ya.M. Paushkin, and Yu. P. Losev, Proceedings of In-

tern. Symp.on

MacromolecularChemistry,Brussels,

1967,121-124,(1968).

39. Ya.M. Paushkin, and Yu. P. Losev, New Petrochemi-

cal Products and Processes, Central Research Institute for PetrochemicalTechnology,Moscow,pp. 50-54 (1971).

40. Ya.M. Paushkin, Yu. P. Losev, and M. E. Elyamberg, J . High-Molecular Compounds, I X , 362-365 (1967). 41. Ya. M. Paushkin, Yu. P. Losev, E. I. Karakozova, and

V.N. Isakovich, High-Molecular Compounds, A, XV,

2496-2500 (1973).

42. E. I. Evdokimov, I. V. Konoval,Yu. I. Firsov, E. A. Vasilenko, A. S. Glebko, V. N. Chkalova,T. L.Zinevich, and A. N. Kopchenkov, Plastmassy (Plastics), 29-31 (1972). 43. U.S. patent no. 4028332. 44. Author’s certificate no. 572472(USSR). 45. Author’s certificate no. 606866 (USSR). 46. Ya.M. Paushkin, Yu. P. Losev, and D.M. B~il’,New

Petrochemical Productsandprocesses, Central Research Institute for PetrochemicalTechnology,Moscow,pp 38-45 (1971).

92

Yu.P. Losev, Ya.M. Paushkin,and V. N.Isakovich, J . High-Molecular Compounds, A. 2502-2505 (1974). 48. Yu.P. Losev, Ya. M. Paushkin, V. N. Maksimenko, E. p. Natal'ina,and v. N.Isakovich, Reports Belarussian Ac. Sci., v. X V I , 916-918 (1972). 49. Yu. P. Losev, V. I. Shonorov, S. P. Baranov, T. S., Mar-

47.

Losev tynenko, E. S. Savostenko, V. N.Isakovich,and V. I. Fursikov, USSR Author's Certificate, no. 63741 1, Invenfor's Bulkfin, no. 46(1978). 50. Yu. P. Losev, Ya. M. Paushkin,and V. N.Isakovich, Abstracts at Intern. Symp. on Destruction andStabilization of Polymers, 257-258 (1974).

7 Poly(malic Acid) from Natural Sources Eggehard Holler

University of Regensburg, Regensburg, Germany

1.

INTRODUCTION

Poly(L-malic acid) denotes a family ofpolyesters derived from L-malic acid as the buildingunit. By chemical synthesis, three kinds of poly(L-malic acid) have been obtained, depending on the molecular position of the ester bond: the a-type(1) [l], the p-type(I1) [2], and the a,p-mixed-type(II1) [3].

p-Poly(L-malate)is a "young"polymer. Because of some of its chemistry it is occasionally counted among owns a the poly(hydroxya1kanoates). However, it wealth of properties and structural diversity not shared by any of these polymers.

The following chapters will be devoted to the production of p-poly@-malic acid)or its salt by fermentation, its isolation, and physico-chemical characterization. The biosynthesis, degradation, andpresumed physiological role will be also considered.

In recent years poly(ma1ic acid) has been discovered in several fungal strains, notably inPhysarurnpolycephalum [4] and in Aureobasidium sp. [5,6], and have been analysed as being of the p-type(I1)either linear ( P .polycephalum) or branched (type IV) (Aureobasidium sp.). Poly(ma1ic acid) is of pharmaceutical interest because its chemical derivatives may harbor both tissuespecific homing molecules and therapeutic effectors to be usedfor tissue (tumor)targeting in chemotherapy [2]. Because of its efficient production by fermentation, its biodegradability andnontoxicity, it isalso considered as raw material in the industrial production of detergents, glues, and plastic materials.

II. PURIFICATION OF P-POLY(L-MALATE) FROM FUNGI A. Fungal Producers of p-Poly(L-Malate) The first organism reported to produce poly(L-malic acid) wasPenicillium cyclopium [7]. An amount of 2.693

94

g freeze-dried poly(L-malic acid) from 8.6 liters of crude fungus extract was purified by repeated anion exchange chromatography and 66% acetone (pH 4) precipitation followed by size exclusion chromatography on Sephadex G50 and concersion of the salt to the polymer acid with Amberlite IR-120. (Rohm & Hans). Based on the finding ofL-malic acid after hydrolysis (hydrochloric acid) and other criteria, the polymer wasassumed to be poly@-malic acid). In the year 1989, plasmodia of P. polycephalum were shownto synthesize p-poly@-malate),the identity of which was confirmedby a variety of analytical methods also involving NMR techniques [4,8]. The polymer is isolated from the liquid growth medium of2-d old cultures (grown as “microplasmodia” during shaking). The polymer can be collected in several batches adsorbed to DEAE-cellulose. The combined batches were subjected to purification by repeated chromatography on DEAE-cellulose, alcohol precipitation, and size exclusion chromatography onSephadex G25 fine and lyophylization. Average yields were 100 mg of polymer in 100m1 of culture medium. p-Poly(L-malate)can also be prepared fromthe plasmodia containingat maximum 80 mg of the polymer in 100 g of cells, but this purification was tedious involving acid precipitation (perchloric acid pH2)of protein in the first step. Highly purified ppoly(L-malicacid)devoidofcontaminating salts was prepared by passage of the purified polymer salt over Amberlite IR-120 (H+ form), lyophylization, solubilization in acetone, removal of debris, and evaporation of the solvent. The final polymer wasobtained in colorless translucent sheets. A detailed description of the purification is given in Table 1. During a large-scale screening (Rathberger, Molitoris andHoller, unpublished results), 232 different fungus strains were tested for the production of p-poly(Lmalate) among them 53 species of marine Ascomycetes, 6 species of yeast-like marine Ascomycetes, 9 species of terrestrial Ascomycetes, 13 species of marine Basidomycetes, 7 species of yeast-like marineBasidiomycetes, 44 species of terrestrial Basidiomycetes, 54 species of marine Deuteromycetes, 15 species of yeast-like marine Deuteromycetes, 19 species of terrestrial Deuteromycetes, Dictyostelium discoideum, Dictyosteliurn mucoroides, Physaum polycephalum, and 7 species of terrestrial Phycomycetes (all cultures Glc-NS-Y medium, see Table 2 footnotes). The best producers were Aureobasidium species of the terrestrial Deuteromycetes (4-9 g/ liter). The strongest known producer of p-poly(L-malic acid) has been identified as Aureobasidium sp. providing 61 g of polymer from 1 liter of culture medium [5,61. pPoly(L-malate, Ca2+-salt)of the culture broth was first separated from accompanying bulk pullulan by methanol precipitation. The water-redisolved precipitate was converted to the polymer acid by passage over Amberlite IR-120B (H+-form). Thus, the best to-day producers of

Holler

p-poly(L-malate) are the Aureobasidiae followed by P . polycephalum (0.5-1 .O fliter), several Cladosporium species of marine Deuterornycetes (0.02-0.35 fliter), and Corollospora species of marine Ascomycetes (0-0.02 fliter). A list of the producers is given in Table 9 L.

B. Methods of Purification of P-Poly(L-Malate) from Fungal Producers

The protocol for the purification of P-poly(L-malicacid) in Table 1 has been worked out for P . polycephalum. It allows an efficient removal of contaminating proteins, nucleic acids, and low-molecular mass impurities from the culture medium or from the cellular extracts of this organism. In order to avoid low-molecular mass polymer fractions, the culture medium has to be harvested early during growth, and the purification has to be limited to only the high-ionic strength eluates after chromatography on DEAE-cellulose andto the high-molecular mass fractions after size exclusion chromatography. The mass-average molecular mass valueof 53 kDa in Table 2 is the maximum rangeobtained but preparations of as low as 12-kDa masses witha high value for the polydispersity (in the range of 3.5-4.0) can be obtained. Extremely low mass valueswere found for the culture medium of macroplasmodia beingof the order of 3-5 kDa. A decrease in molecular mass reflects spontaneous hydrolysis, introducing mostlyintrachain cuts, and the hydrolytic action of polymalatase, a specific enzyme degrading the polymer from one of its ends to L-malate (see the following). Spontaneous hydrolysis is promoted by the slightly acid pH 4.25-4.53 ofthe culture medium during the first 3 days of plasmodial growth. The purification of p-poly@-malicacid)from Aureobasiae has been reported involving methanolprecipitation ofthe polymer in the form of the Ca2+ salt [5]. This is possible because a high concentration of CaC03 is present in the growth medium. Unfortunately, the polymer acid is not soluble in aceton thus missing an additional purification step. In our hands, purification of ppoly(L-malate) from several Aureobasidiae strains was unsatisfactory because of lowyields and resisting impurities. 111.

BIOSYNTHESISOFP-POLY(L-MALATE)

A.

Correlation of Growth and Synthesis of P-Poly(L-Malate) in P. polycephalum

The production of the polymer depends on several factors such as the composition of the growth medium, the time of harvest, and the particular stage of the life-cycle of organism under consideration. For P. polycephalum only plasmodia are the producers of p-poly(L-malate); neither amoebae nor spherules (specialized cell forms that can survive unfavorable environmental conditions)

Poly(ma1ic Acid) from Natural Sources

95

Table 1 Purification of P-Poly(L-Malic Acid) and Its Potassium Salt from the Culture Medium

of Plasmodia ofPhysarum polycephalum 1. Grow microplasmodia in 25 X indented' 2 liter Erlenmeyer flasks, each inoculated with 10-g plasmodia (24 h old) of Physarum polycephalum strain M3CVII in 500 m1 growth medium.* 2. Growth at 21°C (27°C for strain M3CVIII) for 3 d in the dark. 3. Harvest conditioned medium (11 liter) by ~ieving.~ 4. Stirr into 700-g DEAE-cellulose4 during 2 h. 5 . Wash DEAE-cellulose with 10liter of buffer A5/0.3 M KC1 on a Buechner funneL6 6. Elute the polymer with5 liter of buffer A/0.7 M KC1. 7. Polymer-containing fraction7: adjust with 10mM buffer A to 0.35 M KCl, readsorb to 300-g DEAE-cellulose. 8. Wash loaded DEAE-cellulose on Buechner funnel with 3liters of buffer AlO.2 M KCI. 9. Pour column 0 (8 cm). Elute with 1.5 liter of a 0.2-1 M KC1 gradientbuffer A. 10. Dilute active fraction with 2.5 v01 of buffer A. Stir into 500-g DEAE-Sephacel (Pharmacia) (2 h). 11. Wash loaded DEAE-Sephacelon Buechner funnel with3 liter of buffer A/0.2 M KCI. 12. Pour column (08 cm). Elute with 2.5 liter of a 0.2 M KCI-1.5 M KC1 gradienthffer A. 13. Adjust the polymer containing fraction to 70% (vlv) ethanol and keep at minus 20°C overnight. 14. One hour of centrifugation at 16,000 X g to harvest the precipitate. 15. Remove traces of ethanol and dissolve precipitate in a minimum (10 ml) of distilled water. Molecular sieving on a Sephadex G25 fine column (1300 ml,0 5 cm). Assay fraction for P-poly(L-malate)8 and chloride ions.9 Use only saltfree fractions. 16. Obtain pure P-poly(L-malate) potassium salt by lyophylization. Store in the freezer. 17. Obtain free P-poly(L-malic acid) after passage over Amberlite IR 120 (H+-form) (20 m1 bed volume/l g of polymer salt). Lyophylize, dissolve powder in acetone, remove insoluble material by centrifugation, and evaporate acetone from the supernatant. Indentation supports aeration during growth ona culture shaker. The culture medium Glc-NS-Hem (9) contained in 1 liter (adjusted with NaOH to pH 4.5 f 0.05): 10 g Bacto-Trypton, 1.5-g yeast extract, l l-g Dglucose monohydrate,3.45-g citric acid, 2-g KH2P04, 0.6-g CaC12.2H20, 0.085-g FeS04.7Hz0, 0.6-g MgS04-4H20,0.084g ZnS04.7H20, and 0.005-g hemin. The solution of hemin (0.5 gfliter in 1% (w/v) NaOH is autoclaved separately. For strain propagation prepare shaking cultures at 27°C in the dark using indentated Erlenmeyer flasks (100 ml medium, 500 m1 flask, rotary shaker). Transfer every 2 days 2% ofthe cultureto fresh medium. Macroplasmodia are grown on agar plates (see below): inoculation 5-ml with microplasmodia (grown for 24 h and pelleted for l min at 2400 X g, resuspended in 1 v01 distilled water). Spherules are activated as follows: Place filter 20 min at 120°C autoclaved. After coolingto 50°C pour into paper with spherules on agar (two-fold diluted culture medium, 2% agar agar, petri dishes, dry them over night in the dark, incubate 4-5 days at 27°C in the dark, allowing a macroplasmodiumto grow. Scrape the cell from the agar and transfer it to 100-m1culture medium (500 ml indentated Erlenmeyer flask). The cells may be extracted for cellular P-poly(L-malate). Whatman DE 52, equilibrated in buffer A (footnote no. 5). Buffer A contains 10 mM potassium phosphate pH 7. Steps 4-15 are canied out in the cold room at 4°C. 'Assay for P-poly(L-malate) see Fischeret al. 1989 [4]. * The assay for ester bonds is more convenientat this stage [4,10]. Nephelometric assay for chloride ions withAgNO3 under weakly HN03-acidic conditions.

nor spores contain the polymer. The polymer synthetic activity ofthe plasmodia increases together with the cellular mass as a function of the D-glucose concentration in the culture medium (Fig. la). At a concentration of 11g Dglucosefliter routinely used the in growth medium, plasmodia growfor a periodof 4-6 days under a concomitant increase in the polymer concentration in the culture medium (Fig. Ib). Atthe same timethe polymer content withintheplasmodia stays approximatelyconstant. Polymer productionceases during the arrestof growth. From thereonthe polymer content (and the polymer average molecular mass) decreases due to the hydrolytic decomposition to L-malate (Fig.lb). Since the decomposition begins from the onset of growth, it is advantageous

to harvest plasmodia at an early stage of the culture if a polymer of a high-average molecular mass is desired. The decline in polymer content of the culture medium during prolonged times is also observed with other organisms suchas Corollospora fusca M-214 and probably refers there as well to theactivity of specific hydrolases (see polymalatase from P . polycephalum). B.

Biochemistry of P-Poly(L-Malate) Synthesis

The biochemistryofP-poly(L-malate)synthesis has been investigated for P . polycephalum but is far from

Holler

96 Table 2 Producers of P-Poly(~-Malate), Growth Conditions, Amount of Produced Polymer,

and Selected Propertiesof Polymer' of

Molatity Solubility Polymer Growth medium? content' Polymer growth, of days Organism (temperatureTC) medium)

KC1 (mgAiter

5.7

Penicillium6 bran wheat cyclopium Glc-NS-Hem? Physarum polycephalum7 M3CVII18 mediumCulture Cytoplasm"

from elution for of free content4 acid in (mdgacetone cells) (overall content)

M,,

DEAE-cellulose,

(kDa)'

103

X

-

M,

Polydispersity (kDa)

5.0

3 (27)

-

1 X 103

(0.063) 0.35

yes

Yes

macro(micro)plasmodium

0.35 (0.19)

Nuclear extract

Yes

macro(micro)plasmodium

C0.3 0.5 0.6 0.7

(100%) (9%)

1.O

(33%) (54%) (4%)

0.5 0.6 0.7

(20%) (70%)

1 .o

(2%)

(8%)

25'O 4.9

50 7.4

2.0

11

14

29 100 6.5 20 52 92

1.2 1.7 3.7 1.3 1.4 l .7 3.2

17 27 4.8 14

31 29

1 .S

Aureobasidium sp.

A-91t2

GIc-NS,'~(25) 7

61 x lo3

Glc-NS-Y," 20

9.2 x lo3 8.0 X 103

9

insoluble

Aureobasidium pullulans'4

M- 156 T-207

(21)

0.05-0.15 0.3

slightly slightly

0.2-0.5 0.5

slightly

0.6

(70%)

(75%)

4.0 5.0

4.5 5.5

1.1 1.1

(WO)

6.1

6.8

1.1

Corollospora fuscaI4

M-214

Glc-NS-Y, 8 22 (21)

Cladosporium clado~porioides'~

M-203

Glc-NS-Y, 24 (21)7

92 72

M-204

TS-03 Cladosporium herbarurnl4

M-202 M-205

Glc-NS-Y, 56 (21)7

350

' Unless mentioned otherwise, the properties listed are for thepolymer isolated from the growth medium. Organisms are grown in the dark. After saponificationof the polymer, L-malalate is assayed accordingto [4,10]. It may be necessary to purify the polymer on DEAEcellulose before performing the polymer assay. Measurement by gel permeation chromatographyin 0.2 M phosphate buffer pH 7.0 with polystyrenesulfate or polyethylene glycols ([S] in the case ofaureobasidium sp. A-91) as molecular weight standards. Data processingas described in Ref. [ I l l . Ref. [I]. Ref. [4]. * Similar results were obtainedwith Physarum polycephalum strainMICVII. See Table 1 footnotes. lo The molecular masses may be smaller depending on the time of harvest and on the selectionof chromatographic fractions during purification. I t After cell fractionation accordingto Ref. [12]. l2 Ref. [S]. 8% Glucose, 0.3% ammonium succinate, 0.2% succinic acid, 0.04% K2C03. 0.01% KHzPO4, 0.01% MgSO4 X 7Hz0, S ppm ZnSO4 X 7H20, 0.05% corn steep liquor, and 2% CaCO, in deionized water [S]. I4 From the collection Molitoris, Regensburg, Germany. Is 12% Glucose, 0.1% NaNO3, 0.01% KHzPO4, 0.05% KCI, 0.02%MgSO4 X 7H20, 3% CaCO3, 0.01% yeast extract (Rathberger, Molitoris, HoUer unpublished results).

being understood [13, 14, and unpublished results]. As indicated, of the various forms of cells in the life cycle (amoeba, plasmodium, spherule, spore) only the plasmodium produces the polymer. Newlysynthesizedpolymer appears first in cell nuclei, which, however, main-

tain a constant level of the polymer over time (homeostasis). Homeostasis is accomplished by a high rate of polymer synthesis and a release of surplus polymer viathe cytosol into the growth medium.The molecular massdistributionof P-poly(L-malate) during growth

Poly(malic Acid) from Natural Sources

97

D

9 a

200

Y

150

3

100

50

0

1

2

3

4

5

6

7

8

9

1

0

1

1

1

2

7

8

9

101112

D-Glucose [mg/ml] n

I 5 0 0

0

U

1

2

3

4

5

6

Growth time [days]

Figure 1 p-Poly(L-malate) released into the culture medium during plasmodial growth of P.polycephalzrm. (a) Effect of the nutrient (D-glucose) on the growth of plasmodia (--&), and the production of p-poly(L-malate) (--0--). Growth conditions are otherwise as indicated in Table 1 footnote 2. (b) Content of p-poly(L-malate) (+) and Lmalate (--0--) in the culture medium during growth of strain M3CVIII under conditions indicated in Table 1 footnote 2. Inoculation on day 0. Growth termination on day 4.

cally synthesized P-L-[''C]malyl-CoA [ 161 was not used as a substrate and radioactivity was not incorporated into nascent p-poly(L-malate) (Willibald and Holler, unpublished results). Also the desulfoanalogue of coenzyme-A had no effect on the cellular synthesis when injected into macroplasmodia. Homology PCR screening with consensus sequences of bacterial poly(P-hydroxybutyric acid) (PHB) synthases and cDNA from P . polycephalum did not show homologous DNA fragments (Windisch, unpublished results). We think that unlike the bacterial synthesis of poly(P-hydroxybutyric acid), the synthesis of P-poly(L-malic acid) does not involve P-L-malyl-CoA. It may rather use adenylate formation for malic acid activation, which is accompanied by the release of inorganic pyrophosphate from ATP as the second substrate. In vivo injection experiments indicate an inhibition of the polymer production by the ATP analogue a,P-methyleneadenosine 5'-triphosphate (AMP-CPP) but not with P,y-methyleneadenosine 5'-triphosphate (AMP-PCP) (unpublished results) in agreement with this assumption. The polymer synthetic activity seems to be regulated by a signal pathway. The GTP analogue p,y-methyleneguanosine 5'-triphosphate (GMP-PCP) inhibits the polymer synthesizing activity suggesting that a G-Protein is involved. Inhibition was also observed after injection of tyrosine kinase inhibitors, suggesting that one of the cascade proteins in the signal pathway should not be phosphorylated (we have reasons to assume that this is not the P-poly(L-malate) synthase itself). During plasmodium disruption that might activate the signal pathway, an extremely rapid inactivation of P -poly(L-malate) synthase is observed.

IV. PHYSICO-CHEMICAL PROPERTIES A.

is such that the nuclei and the cytosol of a plasmodium contain approximately equal amounts of p-polyl Lmalate. p -Poly(L-malate) is produced from D-glucose involving the citric acid cycle for the production of precursors [13]. Plasmodia have a relatively high content in L-malate (0.3-3 mg/g cells). The content of D-glucose in the culture medium can be varied in the range 0-1 1 mg/ml, and the amount of polymer released into the growth medium varies proportionally (Fig. la). In contrast, the content in the cells stays constant due to homeostasis. By using ~-['~C]glucose in the culture medium or by injecting L-[14C]malate, the production of radioactively labeled p-poly(L-malate) can be preparatively accomplished. In the beginning of the investigation it was believed that enzymatic synthesis of p-poly(L-malate) follows a route, which is similar to the synthesis of poly(P-hydroxybutyrate) (for a review see [15]). However, enzymati-

Molecular Mass

Structural information about p-poly(L-malate) is available for the polymers from P . polycephalum [4, 8, and unpublished results] and Aureobasidium sp. [5,6]. These polymers differ substantially by their molecular mass (Table 2). The Aureobasidiurn-P-poly(L-malate) is of molecular mass below 10 kDa, whereas that of P . poZycephalum in the absence of degradation is of the order of at least 50 kDa and may be as high as several hundred kDa. These two classes of p-poly(L-malate) can be also distinguished by their elution position from the anion exchanger DEAE-cellulose and by their solubility as free acids in acetone (Table 2). Furthermore, the Aureobasidium polymer does not inhibit DNA polymerases (see the following). The producers of the two classes of ppoly(L-malate) are also characteristically different. The high molecular mass polymer is made by P . polycephnZum, a fragile slime mold with extremely large multinucleate plasmodial cells, whereas the low molecular weight polymer is made by cell walls embedded mononucleate cells which display hyphedyeast dimorphismus.

Holler

98

B. SpectralProperties

NMR-spectra have been recorded for the free polymer acid, the Na+, and Ca2+ salts (Table 3). A s the spectra for a- and P-poly(L-malic acids) or the copolymers are very similar and are, therefore, difficult to distinguish [6,8], a comparison of the corresponding high-resolution 'H and 13C NMRspectra has been camed out confirming that the polymer from P . polycephalum was indeed of the poly(P-hydroxy acid) structure [ 8 ] . The absence of band fine splitting revealed also that the natural polymer was of extreme enantiomeric purity exceeding that of synthetic P-poly(L-malate). For P-poly(L-malate) from Aureobasidium sp. A-91, the observed peculiar variation in methine-andmethylenepeaksduringhydrolytic cleavage into L-malic acid was interpreted as evidence for chain branching accordingto the structure (IV) (see Introduction) [5,17]. Branching is also thought to be a reason for the absence of a melting point (decomposition above 185OC) and the insolubility in acetone or other organic solvents [5]. In contrast, P-poly(L-malic acid) (mass-average molecular massof 50 kDa) has a melting point of 192"C, shows decomposition above 200°C

(Gassner andHoller, unpublished results) and is readily soluble in acetone. Infrared spectra and the degree of specific rotation show typicalfeatures of the malic acid polyester (Table 3). Ultraviolet absorbance spectra of P-poly(L-malate) from both P . polycephalum and Aureobasidium sp. A91 are similar andare reminiscent of malate itself [4,5]. For a solution of 1.0 mg/ml polymer, absorbance increases from 0.40 units at 230 nm to 10 units at 190 nm wavelength. After saponification and pH-neutralization, the absorbance increased from 8.7 units at 230 nm to 100 units at 190 nm. According to molecular calculations, p-poly(~malic acid) is highly flexible since long-range ordering forces are absent (Hendl, Urbani and Cesaro, unpublished results). The situation issimilar as found for poly(P-hydroxybutyric acid), which can be considered as the reduced formof P-poly(L-malic acid) with methyl groups in the place of the pending carboxylic groups. Due to charge repulsion in the ionized polymer, an extended conformation is preferred in P-poly(L-malate). Nevertheless, the polyanion probably does not display higher-ordered structures in water.

Table 3 Spectral Properties of P-Poly(malic Acid) and Its Salts

Infrared P-Poly@-malic acid) Na salt* (DzO)

(ppm) 'H NMR 2.9 (methylene,

doublet), 5.0 (methine, triplett)

Na saltt (DzO)

Caz+ salts @20)

3.0 (methylene, doublet), 5.2

(methine, triplet)

Free acidt (CD3COCD3)

3.0 (methylene, doublet), 5.4

(methine, triplet)

* Ref. [4]. t Ref. [8]. $ Meso

diads. 0 Ref. [5]. B Isotactic triads.

rotation I3C NMR (ppm) 36.7 (methylene), 72.3 (methine), 172.3 (ester carbonyl), 176.1

(carboxylate)

36.34 (methylene), 71.95 (methine),S 171.98

(ester carbonyl), 175.78 (carboxylate)$ 38.7 (methylene), 74.2 (methine), 174.2 (ester carbonyl), 178.2 (carboxylate) 36.33 (methylene),$ 69.28 (methine),n 169.07 (ester carbonyl),$ 170.03 (carboxylic acid) 38.4 (methylene), 71.9 (methine), 173.1 (ester carbonyl), 174.6 (carboxvlic acid)

Specific (KBr) spectra, Vmax (cm- ) 3000; 2940 (C-H), 1750-1720 (C*,

ester,

[alD

- 17" (C 6.0

HzO)20°C at

carboxylic acid), 1650-1600 (carboxylate)

3400; 2940; 1740; 1600; 16"-17" 1410; 1280; 1180; 1100; 1050

3400; 2940; 1740; 1410; 1180; l050

(C

5.0,

HzO) at 25°C

- 8" (C 6.0,

HzO) at 25°C

99

Poly(mal1c Acid) from Natural Sources

V. CHEMICAL PROPERTIES OF /~-POLY(L-MALATE) A. Assay of P-Poly(L-Malate) A specific assay for p-poly(L-malate) is alkaline hydrolysis of the polymer followedby the enzymatic dehydrogenation of L-malate and measurement of the NAD/H+ specificA340 absorbance [4]. Hydrolysis is carried out at concentrations of 1-20 pg of polymer in 10O-pI samples overnight (12 h) in the presence of 0.45 M NaOH (10 p1 of 5 M NaOH) at 37°C. Samples that contain protein (cell extracts) were deproteinized after acidification to pH 2 with 7 M perchloric acid, and removalof the precipitate bycentrifugation. The supernatant is then subjected to alkaline hydrolysis.This acid precipitation has the drawback that polymer is lost by adhering to the protein precipitate. A better method is the chromatographic removal of protein on a 1 m1 DEAE-cellulose column (sufficient for an extract from 1 g of cells) by washing with5 m1 of 0.2-0.4 M KC1 in 10 mMpotassium phosphate buffer pH 7.0. The polymer is eluted with 3 m1 of buffer containing 1 M KCl. For a low-molecular mass polymersuch as of 4-8 kDa, the salt concentration in the washing buffer should not be higherthan 0.15 M. After hydrolysis, 10O-pl portions are assayed spectrophotometrically by followingthe A340due to formation of NADH/H in the presence of 830p1 of glycidhydra(0.5 M glycine, 0.4M hydrazine stock zine buffer pH 9.0 solution), 65 p1 of NAD+ (40 mM), and 5 p1 of malate dehydrogenase (30 unitdpl in 10 mM TrislHCl pH 7.5) [B]. The L-malate background inthe sample isassayed for a sample, which has not been subjected to alkaline hydrolysis. The absorbance at 340-nm wavelengthis standardized with known amounts of L-malate. The sensitivity of the assay is approximately0.5 pg of polymer. A 10-foldhigher sensitivity can be obtained with the method of Peleget al. [ 191that couples the malate dehydrogenase reaction with the reduction of 3-(4,5-dimethylthiazol-2-yl)-2,5-diphenyltetrazolium-bromide) (MTT) yielding the corresponding formazan dye. This method has been reported to be applicable as a plate-assay for the screening of L-malic acid secreting microorganisms and may be usedfor the assay of p-poly(L-malate)after a proper technical adaptation. Alkaline (and also acidic) ester hydrolysis of ppoly(L-malate) is accompanied by side reactions leading to the formation of fumarate, maleate and/or racemization, especially at elevated temperatures. The above assays thus underestimate the polymer contents due to the formation of small amounts of 2-4% fumarate (unpublished results). This fraction of fumarate increases for the hydrolysis of more concentrated polymer solutions. A less time consumingalternative to the saponification/malate dehydrogenase assay is the albeit less specific measurementof the carboxylicester groups according to Kakac and Vejdelek [lo]. The sample (160 pl) is +

mixed with a solution of 10% (w/v) aqueous hydroxylammonium chloride (160p1) and a solution of 10% (w/v) NaOH (160pl). After 5-10 min at room temperature, a solution of 4 M HCl (160 pl) is added followed by a solution of 5% (w/v) FeCh (160 pl). The absorbance As40 is read after 10 min. One unit of As40 corresponds to 0.29mg k 0.02 mgof p-poly(L-malate, potassium salt).

B. AcidDissociation p-Poly(L-malic acid) ionizes readily inwater giving'rise to a highly soluble polyanion. Thus, a 2% solution of the free acid of the polymer fromAureobasidium sp. A91 showed a pH 2.0 [5]. The ionic constants have been determined to be pKa = 3.6 for the polymer fromAureobasidum sp. A-91[5] andpKa (25°C)= 3.45 forp-poly(Lmalic acid)of M, 24 kDa fromP.polycephalum (Valussi and Cesaro, unpublished results) Thus, the polymer is highly chargedunder physiological conditions (pH 7.0). C. Noncovalent Binding

of Small Cations

Polyvalent cations such as the protonated forms of spermidine [N-(3-amino-propyl)-I ,Cbutanediamine], spermine [N,N'-bis(3-aminopropyl)-1,4-butanediamine], poly(ethylenimine), poly(L-lysine), histones, etc., bind noncovalently to p-poly(L-malate) with high affinities. For spermidine and sperminethe dissociation constants (pH 7.5, 28"C, 50 mM ionic strength) are K D =~0.6~ ~ mM and = 0.06 mM, respectively [4]. The dye tetranitro blue tetrazolium bindsto the polymer ( K D =~ 0.7 ~ ~mM, pH 7.O,2OoC,50 mM ionic strength) with a concomitant increase in absorbance at 275nm and 350nm wavelength, and canthus be usedas an indicator for studying competitive interactions of other ligands with the polymer [13]. By this method, the binding of Ca2+and Mg2+ was found to follow dissociation constants (20"C, pH 7.0,50 mM ionic strength) of 0.11 M and 0.16M, respectively. These salt complexes are rapidly reversible by the addition ofhigh salt in the molar concentration range. It is of pharmacological interest that cis-diamminepolymalatoplatinum(I1)is obtained from a mixture of P-poly(L-malate)and cis-diamminediaquaplatinum(I1). The platinum complex inhibitsthe growth of MCF-7cell cultures similarly as observed with cisplatin (Natarajan and Holler, unpublished results). D.ChemicalDerivatlration j?-Poly(L-Malic Acid)

of

Molecules ofinterest that contain free amino groupscan be coupled in aqueous solution to p-poly(L-malate) as amides using carbodiimides such as the water-soluble l-ethyl-3(3-dimethylaminopropyl)carbodiimide hydrochloride (EDC) [2,12,20,21]. By this method, the molecules are attached randomly. A selective amide bond formation at the carboxylate terminus can be achieved

Holler

100

without a coupling reagent by a direct attack (aminolysis) of the polymer ester bond, thereby shortening the polymer chain.By using a soluble acylchloride, the terminal hydroxyl group can be esterified in aqueous solution. The yields of such reactions in water are low and the molecular mass of the polymer is reduced by spontaneous hydrolysis. Several kinds of reactions may be carriedout in organic solvents such as N-methylpyrrolidone, dioxane, THF, CHZC12 or acetone, in which the acid is (partially) soluble[23. In principle, reporter, effector and targeting moleculescan be singly or collectively attached to p-poly(L-malate) in order to serve pharmaceutical or other purposes. VI. BIOCHEMICAL PROPERTIES p-Poly(L-malate) from P. polycephalum forms specific complexes with DNA polymerasesa,6, and E from the same organism[4,12,13,22,23]. When boundto the polymer these enzymes are completely inhibited. Activity can be recovered by the addition of spermine hydrochloride. Complex formation (inhibition) wasalso observed with DNA polymerasea from p-poly(L-malate)-producing Aureobasiae (unpublished results) but not with DNA polymerasesfromhigher eukaryotes. In contrast, ppoly(L-malate) fromAureobasidiae was unable to form such complexes. Other nuclear proteins like histones and HMG-like proteins bind also to the polymer from P. polycephalum and have the potential to reactivate DNApolymerasesfrominhibitionwith the polymer [4,22]. These proteins have been shown to bind in cellfree systems with dissociationconstants in the sub-pmolar concentration range. In agreement with these findings, multicomponent complexes of high stability have been foundto be constituents of the nuclei inthe plasmodium of P.polycephalum [12]. The spontaneous dissociation of DNA polymerase from nuclear complexes is very slow andoccurs in minutes to hours. The addition of biogenicamines(like putrescine, spermidine, or spermine), andhistones, like high concentrations of salt, to nuclear extracts accelerate the dissociation of these complexes and causes an activation of the DNA polymerases [4]. Because of their small sizes, the amines have an access to the ion pairs of the proteidpolymer complexes and provoke the rapid dissociation. Model calculations have indicated that p-poly(L-malate) displaysa certain degree of isosterism with the phosphodeoxyribose backboneof DNA (and probably with the backbone of RNA) regarding the distance between the negative charges [221. It is, therefore, possible that P-poly(L-malate) mimicks DNA in many its of activities. VII. EVIDENCEFORPHYSIOLOGICAL FUNCTION Evidence for a physiological function of p-poly(L-malate) is availablefor plasmodia of P.polycephalum. The

plasmodium isa giant amorphous cell containing billions of diploid nuclei that develop through the mitotic cycle with high synchrony. Since the polymer is only found in the plasmodium and not in other cell types of the organism, P-poly(L-malate)can be assumed to be essential. The growing plasmodium contains the polymer in its nuclei at high, homeostatically controlledconcentrations (100-150 mg/ml). The polymer forms large complexes of molcular mass in the range of 1-2 x lo6 Da involving nuclear proteins such as histones, DNA polymerases, and HMG-proteins [12]. The present view of the physiological role of p-poly(L-malate) is a function as a mobile matrix andstorage device for certain nuclear proteins among them histones and DNA polymerases. In this role, p-poly(L-malate) provides an equal supply of material to all nuclei that is necessary to maintain the observed synchrony between nuclei. In its role as a complexing agentthe polymer can participate in nucleus organization. Due to its polyvalency it probably also functions as a molecular chaperon of nucleosome assembly [22]. By its property as a molecular bufferthat binds competitively both histones and DNA polymerases, the synthesis of histones becomes coupled with the availability of free active DNA polymerases at the onset of DNA replication. The secretion of the polymer into the growth medium is not only a means to control nuclear p-poly(Lmalate) concentration but also is a prerequisite for sporulation. L-Malate, whichis a co-inductor of sporulation (Renzel and Hildebrandt, personal communication), is obtained by spontaneous andpolymalatase-catalyzed hydrolysis of the secreted polymer. VIII. DEGRADATION A. Spontaneous Ester Hydrolysis P-Poly(L-malate) decomposes spontaneously to L-malate by ester hydrolysis [2,4,5]. Hydrolytic degradation of the polymer sodium salt at pH 7.0 and 37°C results in a random cleavageof the polymer, the molecular mass decreasing by 50% after a period of 10 h [23. The rate of hydrolysis isaccelerated in acidic and alkalinesolutions. This was first noted by changes in the activity of the polymer to inhibit DNA polymerasea of P.polycephalum [4]. The explanation of this phenomenon was that the degradation was slowest between pH 5-9 (Fig. 2) as would be expected if it were acidhase-catalyzed. In choosing a buffer, one should beaware of specific buffer catalysis. We found that the polymer was more stable in phosphate buffer than in Tris/HCl-buffer. Whereas the cleavage of p-poly(L-malate)at neutral pH is at random [21, alkaline hydrolysisreveals characteristic patterns of the cleavage products, which is due to nonrandom chain scission (Fig.3). The phenomenon is explained by an autocatalytic ester hydrolysis. Assuming that one (or both) of the polymer ends bends

Poly(malic Acid) from Natural Sources loo

101

I

50

0

1

2 3

4 S

6

7 8 9 10 1 1 1 2 1 3

PH Figure 2 Stability of P-poly(L-malate) measured by its activity to inhibit purified DNA polymerase a! of P . polycephalum. The relative degree of inhibition is shown (100 rel. units refer to complete inhibition). The DNA polymerase assay was carried out in the presence of 5 &m1 P-poly(~malate) as described [4]. The polymer was preincubated for 7 days at 4°Cin the following buffer solutions (50 mM): KCVHCI (-A-). Citrate (-V-). 2-(N-Morpholino)ethanesulfonic acid, sodium salt (6). Sodium phosphate 10 20 30 40 50 (-C-). N-(2-Hydroxyethyl)piperazine-N’-(2-ethanesulRetention time [min] fonic acid), sodium salt (-A-). N,N-bis(2-Hydroxyethyll). glycine, sodium salt (-V-). Tris/HCI (+), 3-(CycloFigure 3 Reversed-phasechromatographyofproducts hexy1amino)-l-propanesulfonicacid, sodium salt (-W). after alkaline hydrolysis of P-poly(L-malate). Discrete polymer products are formed, which differ in length by several units of L-malate. The absorbance at 220-nm wavelength was measured. (a) P-Poly(L-malate) before hydrolysis. (b) toward the polymer chain while functioningas the cataAfter 10-min incubation in20 mM NaOH at 37°C. (c) After lyst, the observed cleavage patterns are thought to be 15 h in20mMNaOH at 37°C. (d) After 1 h in 500 mM generated. An autocatalytic hydrolysis was also postuNaOH at 100°C.High pressure chromatography (HPLC) lated for degradation under acidic conditions[17]. on Waters reversed-phase Cl*-F-Bondapak. The methanol gradient (in water-trifluoro acetic acid, pH 3.0) was programmed as follows: 0-40 min 0.3-23%, 40-47 min B. Enzymatic Hydrolysis by Polymalatase 23-40%, 47-49 min 40%, 49-54 min 40-0%. (d) Inset: size The stabilityof the polymer ina conditioned culture meexclusion chromatography after 3-min alkaline hydrolysis dium was strongly decreased by the action of an esterat pH 10.2. BioSil SEC 250 column of 300 mm X 7.8 mm ase. The results shown in Fig. 1 indicate that during size, 0.2 M potassium phosphate buffer pH 7.0. prolongedtimes the content inP-poly(L-malate)decreases givingrise to theformation of L-malate. In conis termed “polymalatase.” It removes L-malic acid from trols with fresh medium this degradation was not deone of the two termini of the polymer. Dueto itsactivity tected over a periodof several days, indicating that the in relatively acidic optimum (pH3-5) it is active only the phenomenon was enzyme catalyzed. An esterase was culture medium.The properties ofthe enzyme are sumisolated after ammonium sulfate precipitation of proteins in the conditioned culture medium, followed by repeated marized in Table 4. It has been found that the esterase inhibitors of serin proteasesdo not inactivate polymalahydrophobic interaction chromatography on butyl-Toytase, and that the enzyme contains an essential sulfhyoperl, and size exclusion chromatography on Pharmacia dryl group.It is, thus, distinguished from the “depolymSuperdex 200 (exclusion limitfor proteins 1.3 X lo6 Da) erases” of poly(P-hydroxyalkanoates), that are related [24]. The hydrolase has a molecular mass of and is 68,000 to lipases. Judged by the above and several other critespecific for P-poly(L-malate). Becauseitsofspecificity it

Holler

102

Table 4 Properties of p-Poly(L-Malate) Hydrolase fromP. polycephalurn'

Molecular massZ Specific activity (m~nits/mg)~ K, (mM) Maximum activity,' pH Temperature Inhibitors, residual activity' (%) KCl, 100 mM CaC12, 100 mM ZnCl2, 100 mM FeCh; CoC12; CuC1; CuSO4 (5-100 pM) EDTA, 3 (20) mM 2-Mercaptoethanol 10 mM Glycerol, 15% (50%) (v/v) p-mercuribenzoate, 0.025 (0.1) pglml N-Bromosuccinimide,8 (12) p M Spermine hydrochloride,0.7 (1 .O) mM 0.3 mM Inhibitors of serin proteases' Dgluconolactone, 22 (200) mM Poly(L-aspartate)'O, 10 mM Poly@-glutamate),10 mM Poly(viny1 sulfate), 10 mM Poly(acrylate), 10 mM

68,000, monomeric 18OOO (pH 3.5)4;9000 (pH 4.0)'; 230 (pH 6.W 11(pH 3.5); 3.5 (pH 4.0); 0.11 (pH 6.0) 3.5 27°C 50

15

80 100 80 (40) 80

50 (10) 50 (0) 50 (0) 50 (0) 170 100 50 (0) 100 100 100 100

Ref. [24]

* SDS-polyacrylamid gelelectrophoresis.

Samples containing appropriate buffers were incubated for varying times with varied concentrations of ppoly(L-malate). In 30 min intervals, 100p1 aliquots were removed and assayed for L-malate content according to Ref. [H]. sodium citrate (20 mM). potassium acetate (30 mM). 2-(N-morpholino)ethanesulfonic acid (20 mM). 'The assay as above was used containing 1.5 pg/ml p-poly(L-malic acid) (9.6 mM in terms of malyl residues). The same assay as in footnote 7 was used. In cases of no inhibition, assays were repeatedin the presence of 0.09-1.8 mM polymer (in terms of malyl residues). For composition of the inhibitor cocktail see Ref. [241. lo For molecular masses Ref. [24].

ria, the catalytic mechanism of polymalatase appears to be closer to that of glycosidases than to that of serinesterases. Substantialamounts of polymalatase have been isolated from plasmodial extracts. This may refer to stored enzyme before secretion, because p-poly(L-malate) is not degradedin plasmodia [24]. Several other fungi were found to secrete p-poly@-malate) degradingactivities to L-malic acid (Ratberger, Molitoris and Holler, unpublished results). These enzymes have not yet been purified and characterized. Polymalatase may be useful for the tayloring of ppoly(ma1ic acid) and its derivatives, and for analytical purposes. If the hydrolase is arrested at points of polymer branches or covalently/physically attached ligands, the hydrolase can be used instudies analogous to those known for DNA and exonucleases.

IX. SYNTHETIC ~-POLY(L-MALATE)

Synthetic p-poly(L-malic acid) can beobtained by polymerization of malolactonic acidbenzylester followed by hydrogenolysis of the resultinglinearP-poly(L-malic acid benzylester) [25], and from L-aspartic acid [26], or L-malic acid[27]. The molecular massof the linear products is in the range for the natural polymer. The biochemical reactivity of synthetic p-poly(L-malate)has been indistinguishablefrom the natural polymer [4,22]. Studies with the synthetic polymer have indicated very low levels of toxicity and high immunologicaltolerance [2]. The natural polymer, which shows a higher degree of optical puritythan the synthetic material, may be superior to the synthetic material with regard to properties such as crystallizability or biocompatibility. The extremelylowtoxicity demonstrated for the synthetic

Poly(malic Acid) from Natural Sources

polymer is compatible with L-malic acid beinga natural metabolite.

X. FUTURE p-Poly(L-malic acid) maybe visualized to a certain extent as a member of the large poly(hydroxya1kanoate) family, which is wide-spread in many bacterial strains, and is available in relatively large quantities by fermentation (sustainable feedstocks). According to the instability ofthe ester bond andthe susceptibility of the building units to metabolic degradation, the polymers are environmentally safe raw materials, the more since many of these polyesters can be actively degraded by bacteria and fungi. Production costs can be relatively low, and some of the polyesters are, thus, eligible as raw materialfor manufacturing plastics. As for plastics, p-poly(L-malic acid) is notcompetitive, because its production costs are much too high. PHB and related polymersare water insoluble, limiting an application to nonaquaeous systems. p-Poly@malate) is extremely water-soluble and is, thus, of complementary nature. Aside qualifying as raw material for the manufacture of water-soluble plastics or tissue, the polyanionic nature allows several other applications, some of which probably justify the relatively high production costs. Such applications are the use as precipitating agent, glue or as a surface film or molecular shuttle for carrying covalently bound drugs, cellular targets and tags. This kind of application should be of value in the fields of pharmacology,medicine,and agriculture. Poly(L-malic acid) may also be applicablein the form of blends with poly(hydroxya1kanoates) or other polymers. As more organismsare discovered as producers, a diversity instructure and compositionmay become available that renders this material even more interesting.

REFERENCES 1. T.Ouchi,and A. Fujino, Makromol. Chem. 190: 1523 (1989). 2. C. Braud, and M. Vert, Trends PolymerS c i 3 : 57 (1993). 3. T. Fujino, and T. Ouchi, Polym. Prepr., Jpn. 35: 2330 (1985).

103 4. H. Fischer, S. Erdmann, andE.Holler, Biochemistry 28: 5219 (1989). 5. N.Nagata, T. Nakahara, andT. Tabuchi, Biosci. Biotech. Biochem. 57: 638 (1993). 6. N. Nagata, T. Nakahara, T. Tabuchi, R. Morita, J. R. Brewer, and S. Fujishige, Polym. J . 25: 585 (1993). 7. K. Shimada, K. Matsushima, J. Fukumoto, andT. Yamamoto, Biochem. Biophys. Res. Commun. 35: 619 (1969). 8. S . Cammas, Ph. Guerin, J. P. Girault, E.Holler, Y. Gache, and M. Vert, Macromolecules 26: 4681 (1993). 9. J. H. Daniel, and H. H. Baldwin, Methods in Cell Physiology (D. A. Prescott, ed.) Academic Press,New York, pp. 9-13 (1964). 10. B. Kakhk, and Z. J. Vejdtlek, Handbuch der photometrischen Analyse organischer VerbindungenI :333 (1974). 11. J. F. Johnson, Encyclopedia of Polymer Science andEngineering 3 : 501 (1985). 12. B. Angerer, and E. Holler, Biochemistry, 34: 14741 (1995). 13. E. Holler, B. Angerer, G. Achhammer, S. Miller, and C. Windisch, FEMS Microbiol. Rev. 103: 1 0 9 (1992). 14. C. Windisch, S. Miller, H. Reisner, B. Angerer, G . Achhammer, and E. Holler, Cell Biol. Znternat. Reports 16: 121 1 (1992). 15. A. J. Anderson, and E. Dawes, Microbiol. Rev. 54: 450 (1990). 16. B. Willibald, H. Boves, and E. Holler, Anal. Biochem. 227: 363 (1995). 17. S. Fujishige, R. Morita, and J. R. Brewer, Makromol. Chem., Rapid Commun. 14: 163 (1993). 18. I. Gutmann, and A. W.Wahlefeld, Method. Enzymat. Anal. 2: 1632 (1974). 19. Y. Peleg,J. S. Rokem, andI. Goldberg, FEMS Microbiol. Letters 67: 233 (1989). 20. G. E. Means, and R. E. Feeney, Chemical Modification of Proteins, Holden-Day Inc., San Francisco (1971). 21. H. Yamada, T. Imoto, K. Fujita, K. Okazaki, and M. Motomura, Biochemistry 20: 4836 (1981). 22. E. Holler, G. Achhammer, B. Angerer, B. Gantz, C. Hambach, H. Reisner, B. Seidel, C. Weber, C.Windisch, C. Braud, Ph. Guerin, and M. Vert, Eur. J . Biochem. 206: l(1992). 23. G. Achhammer, A. Winkler, B. Angerer, and E. Holler, Current Genetics 28: 534 (1995). 24. C. Korherr, M. Roth, and E. Holler, Can. J . Microbiol. 41 (Suppl. I): 192 (1995). 25. M. Vert, and R.W. Lenz, Polym. Prepr. (Am. Chem. Soc., Div. Polym. Chem.) 20: 608, 611 (1979) 26. M. Vert, and R. W. Lenz, Polym. Prepr. (Am. Chem. Soc., Div. Polym. Chem.)20: 609 (1979). 27. Ph. Guerin, M. Vert, C. Braud, and R. W. Lenz, Polym. Bull. 14: 187 (1985).


Stabilization of Polyolefins Sahar AI-Malaika Aston University, Birmingham, England

1.

INTRODUCTION

Oxidation of organic materials is a major cause of irreversible deterioration for a large number of substances. It is responsible not onlyfor the loss of physical properties of plastics, rubbers, fibers, and other polymeric substrates, but also for rancidity of foodstuffs, deterioration of hydrocarbon lubricating oils, as well as biological ageing and is also implicated in some diseases. Inhibition of this oxidation process is, therefore, very important and the art of stabilization of organic materials against the effect of molecular oxygenhas evolved over the past 90 years from an entirely empirical basis to a mature science-based technology, which is inplace throughout the polymer industry. Almost all synthetic polymers require stabilization againstthe adverse effect of their processing, fabrication, storage, and the end use environment. The early intuitive trial-and-error approaches in the development of stabilizers arethe roots to the current state-of-the-art stabilization technology of polymers. II. EARLY DEVELOPMENTS

“Aging” or “perishing” of natural rubber was noted long before Hoffman [l] in 1861 drew attention to the role of oxygen inthe deterioration of organic materials. Practical solutions were sought to overcome the problem of loss of properties with aging. Animalfoodstuffs were preserved by smoking; thisprocess almost certainly results in the absorption of phenolic materials by the fatty components of the food. Similarly, additives were selected to prolong the “life” span of natural rubber. By the early part of this century, considerable progress was made toward finding practical solutions to combat the

woblem of oxidation and property deterioration, and a number of chemical compounds including phenols, quinones, and amines were patented and offered as remejies [2-41. Maureau and Dufraisse [5] in the 1920s believed that the primaryrole of “antioxygens” (i.e., antioxidants) wasto deactivate “peroxides” giving inert products. As such this concept failed to provide a suitable mechanism of antioxidant action; their findings, however,provided a phenomenological basis upon which the modem theory of antioxidant action was founded. This was followedlater by the concept of chain reactions involving free radicals and molecular oxygen. Bateman, Gee, Barnard, and others at the British Rubber Producers Research Association [6,7] developed a free radical chain reaction mechanism to explain the autoxidation of rubber which waslater extended to other polymers andhydrocarbon compounds of technological importance [8,9]. Scheme 1 gives the main steps of the free radical chain reaction process involved in polymer oxidation and highlights the important role of hydroperoxides in the autoinitiationreaction, reaction Ib and IC. For most polymers, reaction le is rate determining and hence at normal oxygenpressures, the concentration of peroxyl radical (ROO’) is maximum and termination is favoured by reactions of ROO: reactions If and lg. 111.

ANTIOXIDANT MECHANISMS AND CLASSIFICATIONS

The development of the autoxidation theory, in which the propagating radicals, alkyl, and alkylperoxyl ( R & ROO), and the hydroperoxide (ROOH) are the key intermediates, has therefore led to a comprehensive theory of antioxidantaction; Scheme 2 shows the two major 105

106

Ai-Maiaika Shear

* Rm

Initiation Ro(Ho)

OjRH

ROOH

)c

RO=+ ~d

(W

Rm+ ROH (H,O)

(c)

I

R'+ 02 Propagation

ROOm+RH

RDS

L

ROOH

+

}

Rm

(dl

2 ROOm ____) Inert Products

ROOm+ Rm-

ROOR

2Rm

R-R

(e)

Chain reaction

Q (g)

(h)

Scheme 1 Free radical chain process involved in polymer oxidation.

antioxidant mechanisms:the chain breaking andpreventive inhibition processes. The basis to the chain breaking donor (CB-D) mechanism, which wasthe first antioxidant mechanism to beinvestigated,waslaiddown by the late 1940s [lo-121. Many reducing agents, e.g., hindered phenols and aromatic amines, which reduce the ROO' to hydroperoxide in a CB-D step have already been empirically selected and used for rubbers and by this time also for the newer plastics industry (e.g., Table la, A 0 1-8 and 9-12). The major mechanistic landmarksof the antioxi-

RH

RO.

dant action of hindered phenols are: (1) the importance of the stability of the derivedphenoxy1radical (see Scheme 3,reactions 3b) inorder to prevent chain transfer reactions (see reactions 3c); and (2) the dependence of the activity of hydrogen donor antioxidants on oxygen pressure (see reaction 3d). Oxidizing agents, e.g., quinones, which were shown to be ableto retard oxidation [ 131can function as antioxidants (via a chain breaking acceptor process, CB-A) if they can compete with oxygen for the alkyl radicals (Scheme 4). In the case of polymers, reaction 4a can

RH

+ OH.

BREAKING

1

ROO.

PREVENTIVE

CB-D - Chain Breaking Donor CB-A Chain Breaking Acceptor

PD WA

MD

ROOH

A

R&

MD

- PeroxideDecomposer

- UV Absorber - MetalDeactlvator

hv M+IM+~

Scheme 2 Autoxidation cycles and antioxidant mechanisms.

Stabilization

107

Table 1a Some Commercial Thermal Antioxidants

Commercial or name common Antioxidant (AO) name common

Antioxidant (AO)

or

Chain-Breaking Antioxidants

Peroxide Decomposers 13 PhosclereP312 Ultranox TLP

(AH) ~ H ~ - R R = -H

1 BHT 2 TopanolO

R=

3 Irganox 1076

R = R =

'9C02q8H17

(m@]-P

Mewe -(aZ09m2)4c

Commercial

14Irgafos168

3

4Irganox1010 5 Ethanox330

P

Irganox 1330

R =C,PS802 Irganox16 z%s 17 Irganox PS800

Roc~m2"z,s R =

Me '

6Irganox 3 114 Goodrite 31 14

M = m.R = C,H9

18 Robec 2 bud (dithiocarbamate)

Metal Deactivators 19 IrganoxMD-l024 R=1098 H Irganox W W 3 20

2246 Cyanox 8 tBu CHI

CH3 [@CH=N-NH-!-h

21EastmanOABH

9 a-Tocopherol

B 22 Mark 1475

11 R1 = H, R2 = HN Nonox DPPD

12Vulkanox

PBN

-0

108

AI-Malaika

Table 1b Some Commercial Photoantioxidants (Light Stabilizers)

Antioxidant (AO)

name common

or Commercial

A 0 code

UV-Absorbers

q H 1 7 1

t-Bu.

t-Bu

23

Rl=RZ=

tBu. R 3 S l

Tinuvin 327

24

%-*U.

Rz=CH3,R3=Cl

Tinuvin 326

25

Tinuvin P Tinuvin 120

26

Cyasorb UV1084 Chimasorb N-705

28

Irgastab 2002

29

Nickel dialkyldithiocarbamate (NiDRC)

30

Nickel dialkyldithiophosphate (NiDRP)

31

Nickel dialkylxanthate (NiRX)

32

Tinuvin 770

33

Tinuvin 622

34

Chimasorb 944

35

Rz=RI=H,

0

Cyasorb UV531 Chimasorb 81

CH3.

R3sH

27

t-BU

Nickel Complexes

109

Stabilization of Polyolefins

Scheme 3 Antioxidant mechanism of hindered phenols.

compete with reaction 4b during melt processing or fatiguing due to thehigh concentration of macro alkyl radicals under these conditions, hence low oxygen concentration at the reaction site. Many stable oxidizing radicals andcertain metal ionswere later shown [14,15] to act by this CB-A mechanism. The early work of Kennerly and Patterson [la] on catalytic decomposition of hydroperoxides by sulphurcontaining compounds formedthe basis of the preventive (P) mechanism that complements the chain breaking (CB) process. Preventive antioxidants (sometimes referred to as secondary antioxidants), however, interrupt the second oxidative cycle by preventing or inhibiting the generation of free radicals [171. The most important preventive mechanism is the nonradical hydroperoxide decomposition, PD. Phosphite esters and sulphur-containing compounds, e.g., A 0 13-18, Table la are the most important classes of peroxide decomposers. The simpletrialkyl phosphites (e.g., Table la, A013) decompose hydroperoxides stoichiometrically

(PD-S) to yield phosphates and alcohols, see Scheme 5 reaction a. Sterically hindered aryl phosphites (e.g., A 0 14) have an additional chain breaking activity, i.e. they react with peroxyl and alkoxy1 radicals during their function as antioxidants (reactions 5b and 5c) [18]. Sulphurcompounds, e.g., thiopropionate esters and metal dithiolates (Table la, A 0 16 and 17), decompose hydroperoxides catalytically, i.e., one antioxidant molecule destroys several hydroperoxides through the intermediacy ofsulphur acids [ 19,201. Scheme 6 shows a simplified scheme for the antioxidant mechanism of simple alkyl sulphides. Metal deactivators (MD) act, primarily, byretarding metal-catalyzed oxidation of polymers; they are, therefore, important under conditions where polymers are in contact with metals, e.g., wires and power cables. Metal deactivators are normally polyfunctional metal chelating compounds (e.g., Table la, A 0 19-22) that can chelate with metals and decrease their catalytic activity [21]. UV absorbers (UVA) act by absorbing UV light hence retarding the photolysis of hydroperoxides. Their activity is also associated with hydrogen bonding between the 2-hydroxy group and the carbonyl chromophore [22]. Typical examples are based on 2-hydroxybenzophenones and 2-hydroxybenztriazoles (e.g., Table lb, A 0 23 and AOs 24-26).

R'

I Propagation h Scheme 4 Antioxidantmechanism of quinones.

IV. EFFECTS OFPROCESSING AND ENVIRONMENTAL FACTORS ON OXIDATIVE DETERIORATION OF POLYMERS

Thermoxidative degradation of polymers can occur at all stages of their lifecycle (polymerization,storage, fabrication, weathering)but its effect is most pronounced

110

At-Malaika

p

y (OR),P=O

+ ROH (a)

+ ROOH [ROP(OH)(OR)3]

1' P(OH)(OR)2+ RO' (a')

(0Ph)sP + RO'

CB

[ ROb(OPh)s]

(RO)(OPh),P

+ PhO'

(C)

Inactive -OI products

Scheme 5 Antioxidant mechanism of phosphites.

during conversionprocesses of the polymer to finished products. Polymer fabrication and conversions are normally achieved throughthe use of high shear mixing machinery (e.g., in extrusion, injection moulding, internal mixing, milling, calendering). Oxidative degradation of polymer articles during outdoor weathering is oftenexacerbated by the combined effects of the environment, e.g., sunlight, rain, ozone, temperature, humidity, atmospheric pollutants, and micro-organisms. The prior thermal-oxidative historyof polymers determines, to a large extent, their photoxidative behavior in service. Hydroperoxides formed during processing (and to a lesser extent during manufacturing and storage) are the primary initiators during the early stages of photoxidation, while the derived carbonyl-containing products (Scheme 7) exert deleterious effects during later stages of photoxidation [23-251. The initiating species, e.g., hydroperoxides, and their decomposition products are responsible for the changes in molecular structure and overall molar mass of the polymer that are manifested in practice by the loss of mechanical properties

(e.g., impact,flectural, tensile strengths, elongation) and by changes in the physical properties of the polymer surface (e.g., loss of gloss, reduced transparency, cracking, chalking, yellowing). The extent of oxidative degradation of the macromolecular chain during melt processing and in-service depends ultimately on the nature and structure of the base polymer. Polyolefins exhibit widely different oxidative stabilitiesdue to both chemical and physical effects. Morphological differences [26], for example, are implicated in the greater susceptibility of polypropylene (PP) and high density polyethylene (HDPE) towards photodegradation than for low density polyethylene (LDPE). PPundergoesmainly oxidative chain scission during processing(Scheme 8). Vicinal hydroperoxides are formed during oxidation of the macromolecular (polymer) chain leading to very high efficiency of thermal and photo-initiation.The breakdown of macromolecular hydroperoxidesduringmelt processing, and subsequently in service under the influence of uv light, yields macroalkoxyl.radicals (PPO.); these radicals which are

example of R is: C&2,0COCH2CH2

Scheme 6 Antioxidant mechanism of simple alkyl sulphides.

111


Scheme 7 Polymer hydroperoxidation during processing and further photolysis of derived carbonyl compounds.

key intermediates in the autoxidation cycle are capable (alkyl, alkoxyl, alkylperoxyl)combination processess, of generating new reactive macroalkyl radicals (reaction see reactions 8a and 8c [27]. 5d) viaa p-scission process. In practice, the deleterious effects of processing on PP cause a dramatic reduction of itsmolecularweight as evidenced by a sharp drop in v. THERMAL STAB~L~ZAT~ON its melt viscosity and high photoxidative instability. In OF POLYOLEFINS LDPE, however, chain scission is less important whereascross-linking reactions (associated within-Polyolefins are exposed tothe effects of high temperacreases in both molecular weight and meltviscosity)pretures initially during processing andfabrication and subof podominate: these reactions occur bydifferentradicalsequentlyduringin-service.Thermalstabilization

-

I

0

CHzCCHzI

R" Cross-linking

-

R 1'

CHZCCHf-+ I

0 I

CH2CCH2I

-CH&

I

II

0

R

Scheme 8 Oxidative degradationof PE and PP.

+

'CH,-

AI-Malaika

112

lyolefins against mechano-oxidative degradation during high temperature processing isessential in order to stabilize the polymer melt andto minimize the formation of adverse molecular impurities and defects, which may, contributeto early mechanicalfailure of finished product during service.The choice of antioxidantsfor melt stabilization vary depending onthe oxidizability of the base polymer, the extrusion temperature, and the performance target of the end-use application. The effectivenessof melt processingantioxidants is normally measured bytheir ability to minimize changes in the melt flow index (MFI) of the polymer that occur in their absence. Chain breakingantioxidants are generally used to stabilize the melt in most hydrocarbon polymers. Hindered phenols (CB-D, e.g., AO-1 to 4, Table la) are very effective processing antioxidants for polyolefins, see Table 2. Aromatic amines, on the other hand, have limiteduse because they give rise to highly colored conjugated quinonoidstructures during their antioxidant function. Although hindered phenols do not suffer as much from the problem of discoloring polymers during melt processing, yellowing can occur as a result of intensely colored oxidation products, e.g., stilbene quinone (SQ) from BHT (Scheme 9) [28]. Transformation products of stabilizers formed during melt processing mayexert either or both anti- and/ or pro-oxidant effects. For example, in the case of BHT, peroxydienones, PxD (reactions 9b, b”) lead to pro-oxidant effects, due to the presence of the labile peroxide bonds, whereas quinonoid oxidationproducts, BQ, SQ, and G. (reaction 9 b’, c, d) are antioxidants and are more effective than BHT as melt stabilizers for PP [29]. The quinones are effective CB-A antioxidants and those which are stable in their oxidized and reduced forms (e.g., galvinoxyl, G., and its reduced form, hydrogalvinoxyl, HG) may deactivate both alkyl (CB-A mecha-

Table 2 Melt Stabilizing Efficiency of Antioxidants in PP (processed in an internal mixerat 190°C). Melt flow index (MFI) Measured at 230°C and 2.16 Kg Antioxidant Control (no antioxidant) Irganox 1010 ( A M ) Irganox 1010 (AM) Irganox 1010 (AM) Irganox 1076(AO-3) BHT (AO-1) Irgafos 168 (AO-14) Irgafos 168 (AO-14) Irganox PS 800 (AO-17) a-tocopherol ( A W ) a-tocopherol (A-) Irganox 101O+Irgafos168

Irganox 1010+IrganoxPS 800

96 w/w

0 0.05 0.10 0.20 0.20 0.20

0.05 0.10 0.10 0.05 0.10

0.054.05 0.054.10

m,g/10 min 11.7 5.9

4.6 3.9 3.7 3.5 7.7 7.2 8.4

3.8 3.6 4.3

4.9

nism) and alkylperoxyl (CB-D mechanism) radicals in a redox reaction (reaction 9e and f). Alkyl sulphides, e.g.,Table la, AO-16and17, which are catalytic peroxide decomposers (PD-C) acting through further oxidation to sulphur acids, are also used as antioxidants during processing of polyolefins (Table 2). However, a major disadvantage of these simple sulphides isthat their conversion to sulphur acids (see reaction Scheme 6 ) involves a parallel series of pro-oxidant reactions involvingthe formation of propagating radicals (e.g., reaction 6e). Therefore, such antioxidants, e.g., AO-17, are always used in combination with effective chain breakingdonor antioxidants, Table 2. Similarly, in the case of phosphites, e.g., AO-14, free radical forming reactions (see Scheme 5 , reactions 5a’, 5b) are believed to occur simultaneously with the stoichiometric peroxide decompositionreaction (reaction 5a), hence the normal practice of using phosphites in combination with CB-D antioxidants (e.g., Irganox 1010) for stabilizing polyolefin melts (Table 2). In the case of long-term thermoxidative stability of polyolefins,stabilizerswith high molarmassesand lower volatility (e.g., antioxidants 4 and 5 in Table 1) are potentially more effective than those with lower molar masses (and higher volatility) containingthe same antioxidant function, e.g., BHT (antioxidant 1 in Table 1). Furthermore, peroxide decomposers, e.g., sulphur-containing compounds, enhance the performance of high molar mass phenols under high temperature in service conditions. For example, dialkylsulphides such as antioxidants 16 and 17 (see Table 1) are often usedas peroxide decomposersynergists in the thermoxidative stabilization of polyolefins.

VI.

PHOTOSTABILIZATION OF POLYOLEFINS

The presence of light absorbing impurities, trace level of metals andadventitious species arising from commercial production processes (polymer manufacture, processing and fabrication) renders manycommercialpolymers (e.g., PE and PP) vulnerable to the deleterious effects of their service environment. The outdoor performance of polymers, however, can be greatly enhanced by a suitable choice of stabilizers. In practical applications, the end-use performance is governed by both physical parameters and chemical factors. Other factors which can affect the ultimate photostability of polymers include samplethickness, polymer crystallinity, coatings, and the presence of other additives, e.g., pigments and fillers. The largest marketfor light stabilizers is in polyolefins; for example photostabilizationof commercial PP is essential for outdoor and indoorend-use applications because of its sensitivity to ultraviolet (UV) light. An effective photoantioxidant must satisfy not only the basic chemical and physical requirements mentioned

113


+ R0

Scheme 9 Oxidative transformation products formed during the antioxidant action of BHT.

above, but must also be stable to UV light and to withstand continuous periods of exposure to UV light without being destroyed or effectively transformed intosensitizing products. Chain breaking donor antioxidants, e.g., hindered phenols, are relatively ineffective under photoxidative conditions as they are generally unstable to UV light and some of their oxidative transformation products are photosensitizing. For example, in the case of the commercial hindered phenol BHT (AO-1, Table l), which is a good processing antioxidant, the peroxydienones (PxD, reaction 9b) undergo photoisomerization [281 through a radical intermediate and can act therefore as powerful photoinitiators (see reaction 9b). Similarly, many sulphur containing antioxidants which are very effective peroxidedecomposers (PD) and thermal antioxidants are not effective photoantioxidants [30]. For example, the intermediate sulphoxide that is formed duringthe antioxidant action of thiodipropionate esters, e.g., AO-17 Table la, photodissociates readily to free radicals (reaction Sf)with deleterious effects on their photoantioxidantactivity. However, both the hind-

ered phenols andthese sulphide antioxidants can synergise with UV stabilizers and become much more effective photoantioxidants. Metal containingsulphur complexes including many metal thiolates, such as dithiocarbamates, e.g., AO-30, dithiophosphates,AO-31, xanthates, AO-32, differ from sulphide antioxidantsin that they are generally morestable to UV light. All metal dithiolates are effective processing stabilizers for polyolefins but their activity as thermal and photo antioxidants depends on the nature of the metalion. Transition metalcomplexes(e.g., M=Ni, CO, Cu) are more photostable than group I1 metal complexes, e.g., Zn, hence better overall photostabilizingeffectiveness(Table 3). Like most other metal dithiolates, the primary oxidationproduct formed during the antioxidant functionof NiDRP (AO-31) is the corresponding disulphidethat undergoes further oxidation to give different sulphur acids (the real catalysts for peroxide decomposition), see Scheme 10 [19]. The UV stabilizing action of nickel and iron complexes (e.g., NiDRC and FeDRC) is strongly concentra-

AI-Malaika

114

Table 3 UV-Embrittlement Times (EMT) of PP and

LDPE Films Containing Different Concentrationsof Antioxidants (processed in an internal mixer at 190°C and 150"C, respectively, and exposed toUV light in an accelerated sunlamp-blacklamp UV aging cabinet) Concentration Antioxidant

Controls (noantioxidant)

I

rnomoog XIO'

PP

EMT, h LDPE

I

0

I

I

I

tion dependent; nickel complexesare much more stable to UV light than iron complexes (Table 3). Further, the behavior of the iron complexes (e.g., FeDRC) changes with concentrations; at low concentration (below 0.05%, in polyolefins) the iron complexes photoxidize rapidly eliminating the thiocarbomoylligandandfreeing the metal ion, which is a very powerfulsensitizer of photoxidation throughits reaction with light and hydroperoxides (Scheme 11 reactions e, f).Conversely, at high concentration, FeDRC shows a photo-induction period during which the iron complex behaves as a photoantioxidant due to the oxidation of the sulphur ligand to low molar mass sulphur acids (by a mechanism similar to that of the dithiophosphate discussed above), (see Scheme 10 and Scheme 11 reaction a, b and Table 3) [31]. This dual role of FeDRC (anantioxidant during processing, storage and very early stages of exposure to UV light, and an effective photosensitizer on further exposure to light) issuitablefor the precise time controlled stabilization requiredfor agricultural applications. Such systems, referred to as the "Scott-Gilead'' systems are now being used for agricultural applications in several countries [32]. The UV absorber (UVA) class (e.g.,2-hydroxybenzophenonesand benzotriazoles, see A 0 23-27in Table lb) arestable to UV light and have high extinction coefficients inthe region 330-360 nm. They operate pri-

lo (Sulphur acids)

Scheme 10 The antioxidant mechanism of NiDRP.

115


b\-

Sulphur acids (AntioxidantEffect)

J

igh Concentration

R'COd + Fe (OCOR')2 Fe (0COR')s

+

R0

+ H20

(Photoactivation) LOW

Concentration

Scheme 11 Mechanismof antioxidant action of FeDRC.

marily by absorbing UV light and dissipating it harmlessly as thermal energy, e.g., via an excited state ketoenol tautomerism.It is important to point out that most known UV stabilizers do not act by a single mechanism; e.g., UV531 ( A 0 23, Table Ib) function not only as a UV screen but also as a sacrificial antioxidant removing chaininitiating radicals (e.g., alkoxy1 radicals) via a weak CB-D mechanism. The limited effectiveness of UV53 1 as a UV stabilizerTable 3) isnot due to its photolability but because of its instability towardhydroperoxides andcarbonyl compounds under photoxidative conmore effective ditions. Thus UV531 is a much photostabilizer for mildly processed saturated polyolefins than for severely oxidized polyolefins[33]. UV absorbers such as UV53 1 synergize with peroxide decomposers (Table 3). The early recognition of the role of stable nitroxyl free radicals, e.g., 2,2,6,6-tetramethyl-4-oxopiperidine, and their hindered amineprecursors, in polymer stabilizationsoonled to the development of the hindered aminelight stabilizer (HALS) class of photoantioxidants. The first HALS, Tinuvin 770, AO-33, (commercialized in 1974)proved to offer much higher UV-stability to polymers than any conventional UV-stabilizer available at the time such as UV-absorbers, nickel compounds and benzoates, Table 3).

Unlike secondary aromatic amines, aliphatic analogues (HALS) are not effective CB-D antioxidants due to the high N-H bond strength, nor are they good UV absorbers or excited chromophore quenchers. of singlet oxygen, Mechanismsinvolvingquenching complexation with important transition metal ions did not appear important. The real antioxidant action of HALS was foundto be due to the ability of their generated stable nitroxyl radicalto scavenge the alkyl radicals in competition with oxygen.This alone, however, could not account for the very high effectiveness of HALS; nitroxyl radical regeneration mechanism was, therefore, proposed andresearches throughout the 1970s and 1980s produced evidence for its regeneration from both the corresponding hydroxylamine and alkylhydroxylamine [34-371, (Scheme 12)[371. VII. FACTORS CONTROLLING ' ANTIOXIDANT EFFECTIVENESS AND USE OF REACTIVE AND BIOLOGICAL ANTIOXIDANTS The effectiveness of antioxidants depends not only on their intrinsic activity but also on their physical retention in the polymer. Migration of antioxidants into the sur-

AI-Malaika

116

ments. These include the use of large molar mass antioxidants, copolymerizing antioxidant functions inpolymers (during polymersynthesis) and the in-situ grafting of antioxidant functions onto polymer backbones during the fabricationandmanufacture of polymer articles using reactive processing methods. Each of these approaches has advantages and disadvantages. The basic tenet here is that if the antioxidant structure can become fully molecularly dispersed in the polymer matrix, and achieve 100% grafting efficiency while retaining its antioxidant function, then the problems of physical loss, migration, and safety of antioxidants will not be an issue. Grafting of reactive antioxidants on the polymer backbone during processing is an important approach to polymer stabilization.Under reactive processing procedures, very high grafting levels of antioxidants can be achieved, hence virtually eliminating the problems of physical loss and migration of antioxidants. Grafting of antioxidants that contain polymer reactive functions (e.g., Scheme 13, AO-I, 11, 111) on saturated polymers can be facilitated by the incorporation of small concentration of a free radical initiator during processing. However, the efficiencyof the grafting reaction (reaction 13b) was foundto be low (about 20-30%) because of competitivehomopolymerization of the antioxidant, reaction 13c [38].Higher levelsof grafting can be achieved when the polymer melt reaction is carried out in the presence of non-homopolymerizable antioxidants (e.g., bis(2,2,6,6-tetramethylpiperidine-4-y1)maleate [39],or by

>N-H

Scheme 12 Antioxidantmechanismof

HALS.

rounding environment, e.g., leaching out into food constituents by the extracting action of oils and fats, leads not only to premature failure of the polymer article but also to problems associated with healthhazards and toxicological effects. Solutions have been sought for enhancing the permanence of antioxidant in polymers especially when used in applications that require direct contact withaggressive solvent extractive environ-

Examples of the reactive antioxidants A" are:

A'

A' (AO-I)

(AO-11)

A'

Scheme 13 Grafting reactions of some reactive antioxidants.

(AO-III)

117


processing reactive antioxidants with a co-agent (a multifunctional monomer without an antioxidant function) [40]. In these case, the efficiency of the melt-grafting reaction can be increased to over 90% together with high efficiency of stabilization, albeit after achieving a delicate balance between the chemical, physical, and rheological characteristics of the reactants. The advantages in adopting a reactive processing approach to polymer stabilizationinclude high retention of antioxidants in the polymer matrix leading to efficient stabilization, reduced risk of migration of antioxidants into the human environment, cost effectiveness and attractiveness of the option of producing highly modified concentrates for use as conventional additives in the same or other polymers, and, site-specific antioxidant grafts can be targetedfor premium performance of speciality niche products. Another approach to safer stabilization is to use a biological antioxidant such as vitamin E (a-tocopherol is the active form of vitamin E, AO-9, Table la). It is essentially a hindered phenol whichacts asan effective chain breakingdonor antioxidant, donating a hydrogen to ROO to yield a very stable tocopheroxyl radical. aTocopherol isa very effective meltstabilizer in polyolefins that offers high protection to the polymer at very low concentration [41], (Table 2).

VIII. SYNERGISM AND ANTAGONISM A cooperative interaction between two or more antioxidants (or antioxidant function) that leads to an overall antioxidant effectgreater than the sum of the individual effects of each antioxidant is referred to as synergism. Synergism can be achieved in different ways. It may arise from the combined action of two chemically similar antioxidants, e.g., two hindered phenols (homosynergism), or when two different antioxidant functions are present in the same molecule(autosynergism);the latter is exemplified by many commercial antioxidants (e.g., Irgastab 2002, A 0 29 Table lb), which have CB and UVA activity. Synergism can also arise from cooperative effects betweenmechanisticallydifferent classes of antioxidants, e.g., the chain breakingantioxidants and peroxide decomposers (heterosynergism) [42]. For example, the synergism between hindered phenols (CB-D) and phosphites or sulphides (PD) is particularly important in thermal oxidation (Table 2). Similarly, effective synergism is achieved between metaldithiolates (PD) and UV-absorbers (e.g., UV 531), as well as between HALS and UV-absorbers, (Table 3). Anti-synergistic effects, on the other hand, arise when antioxidants show antagonistic effects and give rise to a reduced net effect when comparedto the sum of their individual effects [42]. Antagonism during photoxidation of PP occurs when phenolicantioxidants, e.g.,

Irg 1076,AO-3 (CB),are used in combination with metal dithiolates, e.g., NiDEC, AO-30 (PD), due to the sensitized photoxidationof dithiolates by the oxidation products of phenols, particularly stilbenequinones (SQ, see reaction 9C) (Table 3). Hindered piperidines exhibit a complex behavior when present in combination with other antioxidants and stabilizers; they have to be oxidized initially to the corresponding nitroxyl radical before becoming effective. Consequently, both CB-D and PD antioxidants, which remove alkyl peroxyl radicals and hydroperoxides, respectively, antagonise the UV stabilizingaction of this class of compounds (e.g., Table 3, NiDEC Tin 770). However, since the hindered piperidines themselves are neither melt- nor heat-stabilizers for polymers, they have to be used with conventional antioxidants and stabilizers.

+

REFERENCES 1. A. W. Hoffman, J . Chem, Soc., 13: 87(1861). 2. US Patent 99, 935 (1870). 3. US Patent, 680, 387 (1901). 4. German Patent 221, 310 (1908). 5. C. Moureu,and C. Dufraisse, Bull. Soc. Chim., 31(4), 1152 (1922). 6. J. L. Bolland and G. Gee, Trans. Faraday Soc., 42: 236, 244 (1946); J. L. Bolland, Quart. Rev., 3 : 1(1949). 7. L. Bateman,and G. Gee, Proc. Roy. Soc. {a} 195,376 (1948-9); L. Bateman, and A. L. Moms, Trans. Faraday Soc., 49: 1026 (1953). 8. A. V. Tobolsky, D. J. Metz, and R. B. Mesrobian, J . Amer. Chem. Soc., 72: 1942(1950). 9. J. R. Shelton, Rubb. Chem. Tech., 30: 1251 (1957). 10. H. L. J. Backstrom, J . Am. Chem. Soc., 49: 1460 (1927) & 51: 90 (1929). 11. C. D. Lowry, G. Egloff, J.G. Morrell, and C. G. Dryer, Ind. Eng. Chem, 25:804 (1933). 12. J. L. Bolland and P. Ten Have,Trans Faraday Soc., 43: 201 (1947);J. L. Bolland andP.Ten Have,Discuss Farad Soc., 2 : 252 (1947). 13. W. F. Watson, Trans. ZRZ29: 32 (1953). 14. G. Scott, Brit. Polym. J., 3: 24 (1971). 15. G. Scott,Developments in Polymer Stabilization-7,Elsevier App Sci, London, p. 65 (1984). 16. G. W. Kennedy and W. L. Patterson, Ind. Eng. Chem., 48: 1917(1956). 17. S. AI-Malaika, Atmospheric Oxidation and Antioxidants, 18.

19. 20. 21. 22.

vol. I , (G. Scott, ed.), Elsevier Science Publishers, Amsterdam, Chap. 5 (1993). K. Schwetlick, Mechanisms ofPolymer Degradation and Stabilization (G. Scott, ed.), Elsevier Science Publishers, New York, Chap. 2, (1990). S. AI-Malaika, Mechanisms of Polymer Degradationand Stabilization (G. Scott, ed.),Elsevier Science Publishers, New York, Chap. 3 (1990). S. Al-Malaika, K. B. Chakraborty andG. Scott,Developments in Polymer Stabilizationd, (G. Scott, ed.), App. Sci. Pub., London, Chap. 3 (1983). H. Muller, PlasticsAdditivesHandbook, (R. Gachter, and H. Muller, eds.), Hanser, Munich, Chap. 2 (1987). F. Gugumus, Plastics Additives Handbook, (R. Gachter, and H. Muller, eds.), Hanser, Munich, p. 128 (1987).

118

AI-Malaika

23. K. B. Chakraborty and G. Scott, Europ. Polym. J., 15: 731(1977). 24. G . Scott, Developments in Polymer Degradation-l, (N. Grassie, ed.), App. Sci. Pub., London, Chap. 7 (1978). 25. D. J. Carlsson, A. Graton andD.M. Wiles, Developments

in PolymerStabilisation-l, (G. Scott, ed.), App. Sci. Pub., London, Chap. 7 (1979). 26. W.L. Hawkins, Polymer Stabilization, (W.L. Hawkins, ed.), Wiley Interscience, London and New York, p. 7,

(1972). 27. S. Al-Malaika, and G. Scott,Degradation and StabilisaS. Allen, ed.), App. Sci. Publ., tion of Polyolefins, London, Chap. 6 (1983). 28. J. Pospisil, Developments in PolymerStabilisation, vol. I, (G. Scott, ed.), App. Sci. Pub., London, Chap. 1 (1979). 29. T. J. Henman, Developments in Polymer Stabilisation vol. I, (G. Scott, ed.), App. Sci. Pub., London, Chap. 2 (1979). 30. J. R. Shelton, and K. E. Davis, J . Sulfur Chem., 8 : 217 (1973). 31. S. Al-Malaika, A. Marogi, and G. Scott, J . App. Polym. Sci., 31: 685(1986); S. Al-Malaika, A. Marogi, and G. Scott, J . App. Polym. Sci., 33 : 1455 (1987),S . AI-Malaika, A. Marogi, and G . Scott, J . App. Polym. Sci., 34: 2673 (1987).

(N.

32. G. Gilead, Degradable Polymers, (G. Scott, and D. Gi10 lead, eds.)ChapmanandHall,NewYork,Chap. (1995). 33. S. Al-Malaika, and G. Scott, Europ. Polym. J . , 19: 241 (1983). 34. V. YaShlyapintokh, andV. B. Ivanov, Developments in 35. 36. 37. 38. 39. 40. 41. 42.

Polymer Stabilisation-5, (G. Scott, ed.), Elsevier App. Sci., London Chap. 3 (1982). D. W. Grattan, D. J. Carlsson, and D.M. Wiles, Polym. Deg. Stab., l : 69 (1979). G . Scott, Developments in Polymer Stabilisation-7, (G. Scott, ed.), Elsevier App. Sci., London Chap. 2 (1985). S. Al-Malaika, E. 0. Omikorede, and G. Scott, J . Appl. Polym. Sci., 33: 703 (1987). S. Al-Malaika, G. Scott and B. Wyosentono, Polym. Deg. & Stab, 40: 233 (1993). S. Al-Malaika, A.Q . Ibrahim, andS. Al-Malaika, Polym. Deg. Stab., 22: 233 (1988). S. Al-Malaika, and G. Scott, Patent Application Number, PCTiW090/01506 (1990). S. Al-Malaika and S. Issenhuth, Advances in Chemistry Series, ACS, Washington, (R.L. Clough, K. T. Gillen, and N. C. Billingham, eds.) in press. G. Scott, Atmospheric Oxidation and Antioxidants, vol. 2, ElsevierSciencePublishers,Amsterdam,Chap. 9 (1993).

Gamma Radiation Induced Preparation of Polyelectrolytes and Its Use for Treatment of Waste Water T. Siyam Hot Laboratory Centre, Atomic Energy Authority, Cairo, Egypt

1.

INTRODUCTION

Interest in the chemistry of water-soluble polymers (polyelectrolytes) has been continually increasing during the past 45 years. The tremendous scope of utility for water-soluble polymers has ledto a vigorous search for new materials and the rapid development of polyelectrolytes into a dynamic fieldof industrial research. Growth in this field has been especially rapid since 1960; and today, many companies are engaged in synthesis and applications research on polyelectrolytes that are primarily used in four main marketing areas: water treatment, paper, textiles, and oil recovery [l]. Polyacrylamide gel was also used as soil conditioner [2-41. Polyelectrolytes are classified into three main groups: nonionic,anionic, and cationic depending upon the nature of the residual charge on the polymer in aqueous solution as shown in Table 1. The most verstatile and useful type of polyelectrolytes are the cationic, which are comprised of three classes: ammonium (primary [IO], secondary [IIO], tertiary [IIIO] amines, and quaternaries), sulfonium, and phosphonium quaternaries, as shown in Table 2. The extensive industrial and commercial utilization of water-soluble polymers (polyelectrolytes) in water treatment has been developedbased on the charge along the polymer chains and the resultant water solubility. The use of water-soluble polymers in water treatment has been investigated by several authors [5-261 in the recovery of metals; radioactive isotopes, heavy metals, and harmful inorganic residues. This allows recycling water in the industrial processes and so greatly saves 119

the consumption of water in industry. This is evidently of great importanceto meet the requirements of the population, which is continually increasing as is its water consumption. In addition, there is the need to save water for agricultural use to parry the fear of a food crisis. Water-soluble polymers such as nonionic, anionic, cationic, and amphoteric are described as shown in Fig. 1. Polymeric resins such as poly(acry1amide-acrylic acid)[24,25][cationic resin, PAM-AA],poly(acry1ic acid-diallylethylamine-HC1) [20] [amphoteric resin, pAA-DAEA-HCl], and poly(acry1amide-acrylic acid-diallylamine-HC1 [26] [amphoteric resin. PAM-AA-DAAHCl]andpoly(acry1amide-acrylicacid-diallylethylamine-HC1) [26] [amphotericresin, PAM-AA-DAEA-HCl] were also used in water treatment. II. GAMMARADIATION-INDUCED PREPARATION OF WATER SOLUBLE POLYMERS

A. Preparation of Nonionic Polymer (polyacrylamide)

Polyacrylamide (PAM) and copolymers of acrylamide are used on a large scale in waste water treatment and other industrial applications. All of these reasons show that the production and use of polyacrylamide (PAM) and copolymers of acrylamide are a material objective. Acrylamide is polymerized bythe conventionalfree radical initiators, e.g., peroxides [27,28], redox pairs [29-331, and azo compounds [34].Electro-chemicaliniti-

nic

Siyam

120 Table 1 Polyelectrolytes

Nonionic Polyethers Polyamides Poly(N-vinylheterocyclies)

Amonium Amines Quaternaries

Carboxylic Sulfonic Phosphonic

Phosphonium

ation [35], ultrasonic waves [36], photo-chemical sensitizers [37-401, ultraviolet radiation [41], radio-isotopes radiation [42-441, x-rays [45], gamma radiation[45-481, and accelerated electrons [49] also produce PAM. 111.

KINETICS OF RADIATION-INDUCED POLYMERIZATION OF ACRYLAMIDE IN AQUEOUS SOLUTION

Acrylamide polymerization by radiation proceeds via free radical addition mechanism [37,38,40,45,50]. This involves three major processes, namely, initiation, propagation,andtermination.Apartfrom the manysubprocesses involved in each step at the stationary state the rates of formation and destruction of radicals are equal. The overallrate of polymerization (R,) is so expressed by Chapiro [51] as: RP = k,k,"/2Ri'n [M]

(1)

Table 2 CationicPolyelectrolytes

I Amonium

H

I+

amines-N-R-"-R-"-H (Protonated) I H

IO, . IIO, . IIIO

IO 1110

Quaternary

R

l+ -N-R I R

11. -S+-R Sulfonium

I

R

111. Phosphonium

R

l+

-"-R

I

R

H

I+ I

H I10

H

l+

I

R

where k,, kt, and Ri are the propagation, termination, is the and initiationrate constant, respectively, and [ M ] initial monomerconcentration. In homogeneous media acrylamide is terminated by bimolecular termination [51-531. In this case thedegree of polymerization as defined by Chapiro [51] is:

(m,),

p,, = 2 kpkt"'2R~n [M]

(2)

The advantage of using radiation-induced polymerization is that the polymer is homogeneous andfree from any impurities. Moreover, the molecular weight of the formed polymer is controlled by varying doses and doserates.

IV. PREPARATION OF ANIONIC AND CATIONIC POLYACRYLAMIDE Anionic polyacrylamidewas prepared by gamma radiation-initiatedcopolymerization of acrylamidwith sodium acrylate in aqueous solution at optimum conditions for the copolymerization [ 171. The copolymerization process produces water-soluble poly(acry1amide-sodium acrylate [PAM-AANa] of high molecular weight [17,54]. Radical polymerization of diallylamine derivatives producewater-solublepolymers of lowmolecular weight [22,55-571. In order to increase the molecular weight, acrylamide has been copolymerized withthese diallylamine derivatives to produce cationic polymers with variable charge density depending on the content of the structural units of pyrrolidinium rings andacrylamide in polymeric chains [22,55,58-611. Poly(acry1amide-diallylethylamine-HC1) (cationic polyacrylamide PAM-HC1) was prepared by gamma radiation-initiated copolymerizationof acrylamidewith diallylethylamine-HC1 inaqueous solution at the optimum composition for copolymerization of acrylamidewith diallylamine derivatives [61]. A.Kinetics

of Copolymerization Process

1 . Znfluence of Dose-Rate The dose-rate was varied by lining the irradiation chamber of the gamma cellirradiation by lead foil of uniform thickness [17]. The dose was kept constant at 0.15 and 0.35 KGy for copolymerization of AM-AANa and AMDAEA-HCI system, respectively. The results are shown in Figs. 2 and 3, which showthat the rate of polymerization, R, increases while the degree of polymerization and the intrinsic viscosity [ T ] decrease with the dose-rate. The exponents of the dose-rate for AM-AANa system [17,54] were determined to be:

(mII) R,

a

Dp,

Z1.3

0:

I-'.'

(3)

(4)

Gamma Radiation Induced Preparation of Polyelectrolytes (a)

- Non i oni c po I ymers

121

:

CH* "CH2-"

I

c, OH NH2 Polyacry lamide (PAM) "'Structure 1 "

@l-

Anipnit polymers:

I o+!\o-

" C H 2 --CH-""

"

c

0" 'NHZ

c ofi '0-

Na+

polyacr yl amide-sodium,acry late (anionic polyacrylamide pAM-AANa)

Polysodium acrylate(pAAMa1 "5tructu.re 2 "

"

Structure 3"

'iH3

-F

""CH2

= 0

0-

Na+

Polysodium methacrylate ( pAAM Na 1 "

st r uctur e

4

I'

(4- Cat i onic polymers: -c"

- "-

$H3 CH2- C l

COO-CHfCH2

-

y 3 y-HOOC

- CH3

CH3 p o l y #-dimethy laminoet h y l methacry late hydroacytate (POMAEM1 "

Structure 5 "

CH 3 Poly(acry1amide-N,N-dimethylaminoethyl methacrylate (PAM-DMAEM "

+

Na

MC)

Structure 6 "

Figure 1 Structures of water-solublepolymers.

Siyam

122

Poly(dial1 yidimethylammonium choloride-S02 [pDADMAm-Ct-S02] "

Poly(diallyldicthyIammium choloridc-S02 [ pDADEAmCI-S02]

Structure 7"

"

St r uctun 8

"

H \.I

Polyethyleneimine [ * I ] '' St ructure

Poly(N-methyI-2-vinylpyr inium bromide) [ p(2-VPMeBr)I *' Structure 12 '' " "

11 *'

Poly(N-butyl-4-vinylpyridinum

bromide) '*

Ip(4-VPBuBr) St r ucture 13 "

_L

CH2

b-/CH2-"

C H 2 L /

Cl

PoIydiaIIylamine-HCI 1 pMA-HC] "St r udurc

14"

Polydiallylcthylamine-HCI pDAEA-HCI 1 " St ructu re 15 "

Figure 1 (Continued)

1

Gamma Radiation Induced Preparation of Polyelectrolytes

123

Poly(acrylamide-diallylamine-Kl) Po I yamido-amines, pAM- DA A-HCI 1 "Structure 16"

--"CH2-CH

-

I CONH2

CH2-"-

cH2H LL

Poly(acrylamide-diaIIyIdiethylammonium-chloride [ polyami&-amine, pA"DADEAm-Cl ] *' Structure 18 'I

(d)

Amphoteric polymers:

Poly(acrylamide-sodiumacrylate-diallylethylomine-HCI) [ pA"AANa- DAEA-HCI 1 " Structure 19"

Poly(acry1amide-sodium

ocryla

chloride\

e-dial lyidiethylamonium-

[PAM-AANa-DADEAm-CI]

St ructurc 20 *'

124

Siyam

L a,

It will be knownthat for the radical polymerization the increase on the rate of initiation wouldincrease the polymerization rate Eq. (1) and decrease the degree of polymerization Eq. (2). In the present systems, the monomer concentration was relativelyhigh so that initiating radicalsare formed to some extent from the monomer and solvent, i.e., Riin Eq. (1) may be represented as follows [51]:

'11 5

0,

0

R; = f ( 4 d M + ~ s [ S I ) (6) where +M and 4 s are the rate of production of free radiG L 1 cals in the monomer andin the solvent expressed in mole per liter per unit of radiation dose and Z is the dose-rate. Thus, the increase in the dose-rate would increase the rate of initiation, which leads to an increase in the rate 10 15 20 of polymerization, R,, and a decrease in the degree of (1 x lo2 1 , G y l s polymerization, Dp,. The observed increase in the expont of the doseFigure 2 Variation of RP and D, with the dose-rate. rate of the polymerization rate in our systems can be attributed to the peculiarity of acrylamide, which scavenges almost all the radiolytic products of water [63]. The reduction of termination reactions of two propagatHowever, the exponent of the dose-rate for AM-DAEAing radicals, because the reaction becomes highly visHCl system was also determined to be [22]: cous as the polymerization proceeds (trapping of alive radicals) [161. Dainton and Tordoff [401 found that the R, O: P.9 (5) k,lk,0.5 for acrylamide exceeds that of any other polyThe reported values for the exponent of the dosemerization system. Electron and charge transfer prorate for the polymerization rate in gamma radiation-incesses may also play an effective role in multiple initiaduced copolymerizationof acrylamide with methyl chlo- tion of polymeric radicals. These all together leave the ride salt of N,N-dimethylaminoethyl methacrylate system highlysensitive toward radiation. So, the abnor(DMAEM-MC) inaqueous solution was foundto be 0.8 mally high value of the dose-rate exponent might be ex[16]. However, the dose-rate exponent of the polymerpected [54]. ization rate at a lower dose-rate was foundto be slightly Equation (4) shows that the dose-rate exponent of higher than 0.5 for gamma radiation-induced polymerthe degree of polymerization agrees with the theory Eq. ization of acrylamide in aqueous solution [45,62]. (2). However, the degree polymerization andthe intrinsic viscositydecrease with increasingdose rate is probably due to increased termination reactions caused by the increasing radical population at high dose-rate [223.

z

U

I

I

c

2. Influence of Monomer Concentration

IV

a,

'?

c

1

2

3

4

S

6

( 1 x 1 0 2 1 ,Gy I s Figure 3 Effect of dose-rate on the copolymerization of acrylamide with DAEA-HCI. 0 = RP;0 = [v].

Keeping the composition of copolymerizationmedia constant the total comonomer concentration of whichis varied. The absorbed dose was kept constant at 0.14 KGy for the AM-AANa and at 0.35 KGy for the AMDAEA-HCl systems. The results are shown in Figs. 4 and 5 , which show the rate of polymerization, R,, the degree of polymerization, and the intrinsic viscosity increase with increasing monomer coEentration. At comonomer concentration >2. l M L , DP, decreases with increasing comonomer concentration. From the logarithmic plots, exponents of the comonomer concentration for the A"AANa system were determined to be [17,54].

R,

a

IF, O: [ M ] ~ . ~ [M] ~ < 2.1 ML

(7) (8)

nduced on Radiation Gamma

of Polyelectrolytes

s. 0 3.0 U

'S X

1.0

[* r O

0.4

a= l0

, , I 01 2.0 4.0

0 .l

( Comonomers

Figure 4 Variation of RP and concentration.

m,

a

[M]-0.'5

l, M1

L

m,with the comonomers

[M] < 2.1 W L

(9)

In addition, the exponent of the comonomer concentration for AM-DAEA-HCl system was also determined to be [22]: R,,

(10)

D:

The reported values for the exponent of the monomer concentration for the rate of polymerization were found to be 1.26[61],1.3[16] for gammaradiation-inducedcopolymerization of acrylamide with N,N-diethyldiallylammonium chloride and methylchloride salt of N,N-dimethylaminoethyl methacrylate (DMAEMMC). Ishigue and Hamielec [34] have shown that the

125

monomer exponent varies from 1 .O to 1.5 in acrylamide polymerization by chemicalinitiator. In the present systems water is subjected to radiolysis to form radiation products such as: &e; H', OH-, H30+,HZand H202[64]. The three primary radical species react with monomers to give radicals derived from monomers [63]. All these generated radicals are available to contribute to the chain initiation. This increases the exponent of the monomer concentration in these systems. Figures 4 and 5 show that the degree of polymerization and the intrinsic viscosity increase with increasing the comonomer concentration. Thus, theincrease in the comonomerconcentration would increase p,, and, consequently [17] Eq. (2). However, the exponent of the monomer concentrationfor the degree of polymerization decreases at a highcomonomer concentration. The abrupt change in the exponent at comonomer concentration >2.1 M/L (Eq. 9)may be ascribed to polydispersity. This was foundby others [34] to be at 2.1 and 2.2 MIL, although it starts theoretically at 2.0 MIL.

3. Influence of Radiation Dose The influence of the radiation dose on the copolymerization of the AM-AANa system was investigatedat 1.4 MI L [17,54]. The results are shown in Figs. 6 and 7. The data for the copolymerization of the AM-DAEA-HCl system at a comonomer concentration of 1.344 MIL [22] are also shown in Fig. 8. Figures 6 and 8 show, generallythat the conversion first increases fast then slowly with the dose. This is in agreement with the findings of Azzam [48] and Siyam [61]. The significant increase in the conversion percentage isattributed to the gel-effect [51,65]. In the gel-state, as the conversion percentage increases the viscosity of the medium is highly increased and the growing poly-

IO

-

7

B

m \

0

E 5 -

V

F '

L-I

3

U

2

0.5

0.7

1.0

(Comonomers

1

1, M I L

Figure 5 Effectofcomonomerconcentrationoncopolymerization of acrylamide with EAEA-HC1. 0 = RP, 0 = [TI.

A

10'

lo2 Dose , Gy

1o3

Figure 6 Influence of the absorbed dose on the conversion (9). 0 = 0.7035; 0 = 1.407; 0 = 2.1105 M/L.

126

’

Siyam

Table 3 Influence of Radiation Dose on the Copolymer

lo

Composition and the Swelling Degree Dos % KGY degree Acrylamide Amine

W I

2 X

3

IS

24 0.06

0 10’

10

lo2

D o s e , Gy

Figure 7 Influence of the adsorbed dose on the weightaveragemolecularweight. 0 = 0.735; 0 = 1.047; 0 = 2.1105

%

0.10 0.22 0.60 1.12 6.00 14.00 29.00 42.00

21.26 23.52 26.18 28.68 31.25 41.24 42.56 44.41 45.40

78.74 76.48 73.82 71.32 68.75 58.76 57.44 56.59 54.60

Swelling Soluble Soluble Soluble Soluble Soluble 542 750 360 311

M/L.

meric chains are trapped in the viscous medium. Thus, the termination by mutual interaction of two growing chains becomes very unlikely while propagation continues. The interactions of high molecular weight growing chainsbecomediffusion-controlled reactions [66,67]. Consequently, the molecular weight and the intrinsic viscosity of the formed copolymer would increase, as shown .in Figs. 7 and 8 for AM-AANa and for AMDAEA-HCl systems, respectively. The influence of radiation dose on the polymer composition and the swelling degree of (PAM-DAEA-HCl) are shown in Table3.The results show that the percent of acrylamide in the copolymer is higherthan that of the amine. This can be attributed to smaller reactivity ratios of monomers of diallylammonium salts relative to acryl-

amidemonomer [59,60]. Atlow doses water-soluble polymers are produced, the intrinsic viscosity of the resulting copolymers are shown in Fig. 8. At high dose, >6 kGy, the formed polymer converted is into polymeric gel. The gel formation isattributed to polymer branching [51], plurimolecular aggregates[68], and supermolecular architecture [52]. Table 3 also shows that the swelling degree increases then decreases when increasing the exposure dose. The decrease in the value of swelling degree is probably due to the increase in the extent ‘of crosslinking betweenthe polymeric chains as a result of radiation. It was found that the polymerization of monomer of diallyammonium salts proceed byintra- and intermolecular cyclization to formpolypyrrolidinepolymers [64-771, which contain pyrrolidinium rings alternating along the polymeric chains [61]. Consequently, the copolymers includethe amine as pyrrolidinium rings with acrylamide unitsalternating along the polymeric chains dependingon the composition of the copolymer, as shown in “Structure 21.”

10.0

c2ki

2.5

K)’

to2

lo3

0.0

-.-

lo4

Polymerization time,min Figure 8 Effect of radiation doseon copolymerization of acrylamide with DAEA-HCI. 0 = q(%); 0 = [TI.

“Structure 2 1 ”

4. Influence of Temperature The influence of temperature on the copolymerization was investigated at constant absorbed dose of 0.12 and 0.16 KGy for copolymerization of AM-AANa [17,54] and AM-DAEA-HCl [22], respectively. The results are shown in Figs. 9 and 10, which show that the R, values increase while the intrinsic viscosity and the degree of polymerization decrease with increasingthe polymerization temperature. However, the increase in the temperature of the polymerization medium increases the swell-

nduced on Radiation Gamma

of Polyelectrolytes

127

ERp =

n

P

U

Figure 9 Effect of temperature oncopolymerizationof acrylamide with DAEA-HCl. 0 = RP;0 = [TI.

ing of the formed copolymer in the comonomer feed. This increases the concentration of the comonomer around the growing chains as thetemperature rises. So, the propagation reactions and the polymerization rate should increase [78,79]. The increase in the temperature reduces the viscosity of the polymerization medium which increases the termination reactions. This isattributed to an increase in chain transfer reactions higher than that of propagation reactions [16,511. Consequently, the weight-average molecular weightof the formed polymer decreases. From the logarithmic plotof the Arrhenius equation shown in Figs. 8 and 9, the overall activation energy, ER^, was calculated to be 0.65 and 0.56 Kcal/mol for AM-AANa and AM-DAEA-HCl systems, respectively. However, the corresponding reported values for gamma radiation induced copolymerization of acrylamide with DMAEM-MC in aqueous solution was found to be 2.0 KcaYmol [161. In general, the overall activation energy for the polymerization rate is given by:

p+

('Ep -

$)]

where, Ei, E,,, and Er are the activation energy of the initiation, propagation, and termination, respectively. In radiation-induced polymerizationEi is considered to be approximately 0. The Ep value for the radical polymerization is known to be 6-8 Kcal/mol [16,51]. The determined values of ER indicate that Er is about 10.7-14.7 and 10.88-14.88 &al/mol for AM-AANaand AMDAEA-HCl systems, respectively. These large values of Er are presumably due to the diffusion barrier of propagation chains and to the gel-effect.

5. Influence of After Effect The influence of the after effect was studied at an absorbed dose of 0.12 KGy and a temperature of 0°C for copolymerization of the AM-AANa system [17,54]. The results are shown in Fig. 11. Figure 11 shows that the conversion and weightaverage molecular weight increases after ceasing irradiation. The copolymer yield seeminglycome to anequilibrium within6 h, the molecular weightcontinues increasing for nearly 50 h. These may be arguedto be structural changes [80], possibly viaconformationalequilibria [81]. Aggregation [82] and cluster formation [52] may not be neglected. Infact, there is ampleevidence on segmental motion of these and related copolymers [83,84]. V. USE OF WATER-SOLUBLE POLYMERS IN WATER TREATMENT Based on the application of the established theory of colloid stability of water treatment particles [8,85-881, the colloidal particles in untreated water are attached to one another by van der Waals forces and, therefore, always tend to aggregate unless kept apart by electrostatic repulsion forces arising fromthe presence of electrical charges on the particles. The aggregation process

-

12

h

P C

1.00-

.-0

t? Q, > C

8

X

0.60

a2 a20 0 0

12 12

24 24

36 36

48 48

0

Timein hours Figure 10 Variation of RP and

m,,with temperature.

Figure l 1 Influence of after-effect on acrylamide-acrylate copolymerization.

128

is aided by the addition of simple electrolytes, which reduce the electrostatic repulsion. This iscommonly called “coagulation.” Aggregation can also be brought on by various polymeric substances, which bind particles together with the chain molecules forming interparticle bridges. Following la Mer’s suggestion [89,90], the term flocculation,refers specificaly to the latter process. An understanding of the action of polyelectrolytes in flocculation has evolved froma number of investigations of the interaction of various polymeric flocculants [6-20,91-100]with suspended particles. The bridging theory of flocculation was first postulated by Ruehrewein and Ward [5].This theory postulates that the polymer flocculation can be analyzedin terms of a primary adsorption process followed by a second flocculation process [6,8,10,11,14,15].The polymer moleculesattach themselves to the surface of the suspended particles at on or more adsorption sites, and that part of the chain extends out into the bulk of the solution. When these extended chain segments makecontact with vacant adsorption sites on other particles, bridges are formed. The particles are thus bound into a small floc, which can grow to a size limited by the degree of agitation and the amount of polymer initiallyadsorbed on the particle surfaces. If too many adsorption sites are occupied, bridging will be hindered and wholly inhibitedif all are occupied. If too few sites are occupied, bridging maybe too weak to withstand the shearing forces imposed by even mild agitation. Consequently, in the first process of the adsorption the extended chain segments bind particles or ions together with polymerchains. These chain segments interact with each other to form a floc or a sludge in a subsequent flocculation process. The adsorption of polymer molecule to the adsorbent surface may occur by chemical or physical interaction, depending on the characteristics of the polymer chain andthe adsorbent surface. For nonionic polymers, the polymer isadsorbed via hydrogen bonding [8,11,15]. The adsorption of anionic polymers by the suspended particles is considered to be either hydrogen bonding, anion exchange, or chemical reaction with the charged particles [5-7,10,11,19]. However, cationic polymer is adsorbed on the surface of the particles via cation-exchange mechanism [7,11,19]. It was found that [5-71 the rate of flocculation of particles produced by the bridging action of polymer is the slower process and, consequently, the rate-determining step. The primary adsorption of polymer is fairly rapid, butthe slow attainment of the adsorption equilibrium under agitation arises at least in part from the breakdown of flocs offering new surfaces for adsorption. Thus, the bridging step is slow because a polymer adsorbed onone particle must findanother particle having a free surface available to complete the bridge. The interparticle bridging mechanism was affected by the charge [7],adsorption sites [5] ofparticles extending in the flocculation system, the physical characteris-

Siyam

tics of polymers [7,8,10,11,13,101], the degree of polymer adsorption [7],the valency and charge of ions in solution [12,181, the thickness the nature of adsorbed water layer [14,201, and the degree of agitation [5,11]. Inpreviousworks [18-20,23,102] water-soluble polymers such as: polyacrylamide (PAM), polysodium acrylate (pAA Na), poly(acry1amide-sodium acrylate) (PAM-AA Na), poly(acry1amide-diallyethylamine-hydrochloride) (PAM-DAEA-HCl), and poly(acry1amidesodium acrylate-diallyethylamine-hydrochloride) (PAMAANa-DAEA-HCI)were used in the recovery of cations and some radioactive isotopes from aqueous solutions. It was found that the floc is formed between the added polymer and ionsof the solution in the flocculation process with the formation of a crosslinked structure. The formed cross-linked structure is characterized by [103-105]: 1. The transition bond energy is strong energy. 2. The metalionisreadilyremoved by treating with an acid. 3. The coordination number andstereostructure of the metal complexare specific for a given metal ion species. 4. The crosslinking reaction is detectable without difficulty by spectroscopic and magnetic measurements. The efficiency of flocculation of these polyelectrolytes was investigatedat different pH and polymer concentrations, valency of cations, and weight-averagemolecular weightof the polymer. The results are shown in Figs. 12-16.

n

I

U

L

Figure 12 Effect of polymer concentration the on residual metal concentration [M] for p(AM-AANa). 0 = Mg2+, 0 =

W+.

Gamma Radiation Induced Preparation of Polyelectrolytes

129

100

80 60

-

S 40 I U 20 0 5 15

10

[~otymer], mgm/L Figure 13 Effect of polymer concentrationon the residual metalconcentration[M] for p(AM-DAEA-HCl). 0 = Mg2+,0 = Cu2+. Figure 16 Influence of polymsr concentration and weight-averaggmolecular weight(Mw2_Onthe polymer efficiency. 0 = M, = 5.5 X lo6; 0 = M, = 3.0 X lo6.

A.

0

S

(0

15

[polymer], mgm/L Figure 14 Effect of polymer concentration on the residual. metal concentration [M] for p(AM-AANa-DAEA-HCI). 0 = Mg2+,0 = Cu2+.

I

1

Effect of pH

The influence of pH on the residual metal concentration [M] was studied at a constant polymer concentrationof 10 mg4 and copper sulphate concentration of 10 gm/l, results are shown in Table 4. It is clear that [M] decreases with increasing pH value. Results are given in Table 4, which shows that the concentration of Cu2+ decreases withan increase in the pH value for each polymer used. Thisis attributed to theeffect of the pH value on the active groups, which are distributed along the polymer chains. At a low pH value the amide groups

Table 4 Effect of pH on the Percent of Residual Cu2+

Concentration Amphoteric Cationic Anionic polyacrylamide polyacrylamide (PAM-DAEAHCl) (PAM-AANa) pH

Figure 15 Influence of valency of cations on the polymer efficiency. 0 = Eu2+, 0 = Co2+.

5.25 5.37 5.40 5.50 5.56 5.70 5.80 6.15

%[CU]~+ pH 95.36 82.18 30.25 23.22 46.44 11.82 7.31 00

5.22 5.30 5.40 5.65 5.70 5.90 5.95 6.50

polyacrylamide (PAM-AANaDAEA-HCl)

%[CU]~+ %[CuI2+ pH 95.42 90.51 75.21 50.50 30.40 10.21 00 00

5.29 5.25 5.31 5.42 5.52 5.81 6.10 6.14

80.24 60.15 55.02 36.48 38.52 30.31 5.25 00

Siyam

130

are protonated [13,1051, which can undergo imidization [13], while both the carboxylate groups [7,111 and quaternary ammonium groups [61] are unionized. The polymer chains are oftencoiledin the solution.Consequently, no interaction can occur between the coiled polymeric chains and the ions that might be present in the solution. At higher pHvalues the amide groups are unprotonated [106], while the carboxylate groups [7,11] and quaternary ammoniumgroups [61] are ionized.Alternatively, the interaction can occur between these cations of the solution and the active groups of the polymer chains for floc formation.

B. Effect of Polymer Concentration The influence of polymer concentration on the percent of residual metal concentration [M]was studied at a polymer concentration of 10 mgA, metal sulphate concentration of 10 g d , and pH 6, results are shown in decreases with increasing Figs. 12-14. It is clear that [M] polymerdosage.As the polymer concentration increases, the number of extended segmentsonto the solution increases. These extended segments, which affect the probability of the segmentcollisionand consequently the probability of forming the floc, increase. This isin agreement withthe findings of other investigators [7,10,11,13,171. The change in the polymer dosage required to bring 0% [MI varies from one type of polymer to another. This is due to the difference in the weightaverage molecular weight of the polymer used. Figs. 12-14 show that the polymer concentration for complete precipitation of Mg2+ is higher than that of Cu2+. This is due to the difference in the hydration of cations. Increases in the hydration decrease the polymerefficiency. This is in agreement with the findings of Dolimore et al. [14]. Theyfound that the extent of flocculation decreases with increasing thickness of the adsorbed waterlayer around the cations of the clay particles upon treatment with polymeric flocculants.

-

C.

Influence of Valancy of Cations

Different samplesof aqueous solution containing radionuclides of Co2+ and Eu3+ were prepared at different copper sulphate concentrations and constant polymer concentrations (PAM) of 15 mgA. The addition of salt to the system was done to reduce both the repulsion forces between the radionuclides and the interaction between the polymeric chains [7]. The polymer efficiencyfor the prepared samples wasdetermined, results are shown in Fig. 15. It is clear that the polymer efficiencyfor Eu3+ is higher than for Co2+. This can be explained by the difference in the tightly bound structured water associated withdifferent cationic species [14,107]. On this basis, we expect that Co2+is morehydrated than Eu3+. This is due to the difference inthe ionic size. The hydra-

tion energydecreases with increasing ionicsize. This is in agreement with the finding of Dolimore et al. [14]. They foundthat the extent of flocculation decreases with increasing thickness of the adsorbed water layer around the cations. They also found that the flocculation of Li by polyacrylamide froma kaolinite clay is less than that of K and Na. This can beattributed to the strong hydration of Li than of K and Na. D. Effect of Polymer Concentration and Weight-Average Molecular Weight

The influence of polymer concentration on the polymer efficiency (F) was studied at constant copper sulphate concentration of 600 mg/l at two different MWof 3.9 x lo6 and 5.5 x lo6, results are shown in Fig. 16. It is clear that the polymer efficiency (F) increases with increasing the polymer dosage as discussed above. It is found that at high polymer dosage >l00 mg/l, the floc is not formed. This is due to the high viscosity of the solution andthe complete interaction of polymer chains with each other. It was observed that the polymer eficiency (F)increases with increasing molecular weight the of polymer (Mw).The polymer efficiency(F)reaches 100% at polyof 3.9 X IO6 and mer dosage of 24 and 15 mgm/l at 5.5 x lo6, respectively. Consequently, the higher is the value of MWthe lower the polymer dosage. This is attributed to increasing the efficiency of the polymer with increasing weight-average molecular weight of the polymer. On increasing the extended polymer segments are longer making collision and thus bridging is more probable. This increases the probability of forming the floc. This is in agreement with the findings of other investigators [7,10,11,13,15,17].

aw

aw,

VI.

MECHANISM OF INTERACTION OF WATER-SOLUBLE POLYMERS WITH IONS IN AQUEOUS SOLUTION

The infrared spectra of the formed floc from copper sulphate withwater-solublepolymers such as PAM, pAANa, p(A”AA-HCl), andp(AM-DAEA-HCl) are summarized in Tables 5 and 6. In the spectrum of PAM the shift in the absorption bands, characteristic of both >NH stretching and >NH bending to a lower wavelength, is due to the hydrogen bonding formation between the polymer chains in plurimolecular aggregates of polyacrylamide as shown in “Structure 22” [79,108,109]. The same shift in the absorption bands, characteristic for +H group to a lower wavelength, is due to the presence of water molecule, in addition, the appearance of the absorption bands characteristic of the sulphate ion indicate the presence of the sulphate group in the floc. The absorption bands can lead to the conclusion that: (1) polyacrylamide forms complexes with

Preparation Gamma Induced Radiation

of Polyelectrolytes

131

each other; (3) the sulphate ions may be present in the floc. Copper sulphate form crosslinked structures with PAM as show in the possible Structure 23. This is in

"- /\/H l l I

I

-"$"I - H\

,C=O--

H/Y

""q-CH "

"I-

Structure 22

"

Plur irnolecular aggregates of PAM

Cu(I1) and these complexesmayby an octahederal structure due to formation of coordination bonds between the amide group and water of hydration with Cu(I1); (2) the polymerchains form hydrogen bonds with

x = H20 and

YZ = SO42-

'' Structure 23" Cross-linked structure of pAM-CuS04*5H20 complcms

Table 5 IR Spectral Data of Polymers with Copper Sulphate ~~

Absorption bands of Absorption bands

of original groups (cm")

P(AW

P(uNa)

~~

Amide:

>

Free stret. NH (Amide 11) Bondedstret. NH (Amide 11) bend. >NH stret. 1651 (Amide I)

>c"o

Aliphatic:

3270 3390, 3180 3350, 1590

1620, 1650

Blending ------> granulation A

I

I

I

" " " " " " " " " " " " " " " " " " " " " " " " " " " " "

>

Tpu

I

l V

Quality Control of Products Scheme 1

Huang

142

Table 3 The Mechanical Properties of the 566 TPUPVC Polymeric Blends

TPUPVC (ratio by weight) 0.500 25

0.821 15 1.08 13 1.34 13 2.00 26

Tensile strength (MP4

Ultimate elongation

Permanent deformation

(%l

(%l

22.6

304 423 466 553 720

Hardness (“ShA)

71 64

According to Chinese National Standards GB 1040-79, GB 528, GB2411-80. PVC:TPU:Plasticizer:Additives = m:n:5:4 (by weight, m + n + 5 + 4 = 200).

2. TPU Since TPU is much more expensive than PVC, and its quality is usually not as stable as that of PVC, especially when TPU issynthesized instead of a commercial product, the quality inspection of TPU must be done under strict control. One thingthat needs to be paid attention to is the water content of TPU. TPUs with higher than allowed amounts of water must be preheated to remove the bulk of the water. It is always necessary to check the compatibility between the selected TPU and PVC before large scale production. A simple and practical method is to blend a small sample in the lab with a twin-roller calender or a small extruder. A compatible blend should have good processing properties along with a smooth surface and cross-section. If needed, further tests on mechanical properties should be carried out on testing samples made from 2mm films produced by compression molding. B. TheBlending PVC resin, plasticizers, and additives are weighed according to the recipe and mixed in a high-speed mixer. This mixture then is processed at 140-180°C for a certain period of time to form the preplasticized PVC.The TPU is then added and further blending is performed until completion, depending on each case.

If TPUis the major component, then the TPU should be processed first, then added to the preplasticized PVC to form the final blend. The processing (blending) can be done on either calenders or extruders. Usually, one-step processing is feasible oncalenders, but for two-step processing extruders are necessary. The first step is making semi-plasticized PVC granules by the extruder. Then this is mixed with TPU in the proper ratio and added to the extruder once again for final processing.By using a twin-screw extruder, one-step processing is also possible. Since quite a bit of difference exists between raw materials, the recipe, and the equipment, the processing procedure and conditions vary a lot. Also, the processing procedures of commercial products are usually not available to the public. Thus, much work needs to be done to find the best procedure and condition for each individual system. In general, a good procedure is a combination of optimal processing time, temperature, and rotating speedof the screw (in the case of extruder use) or the roll nip (in the case of calender use). V.THE PROPERTIES OF THE PVC/TPU BLENDS

A. GeneralRemarks In general,the properties of the PVC/TPU blends reflect the ratio of PVC and TPU in the mixture.

Table 4 The Low-Temperature Impact Test 39.3%

Temperature Commercial soft PVC, plasticized 566TPUPVC with TPU weight) (by TPU weight) (“C) (by grade medical DOP, by

-25 - 30 - 40 -60

no pass no pass

-

-

According to Chinese National Standards GB5470-85.

-

pass pass

566TPUPVC with

-

54.7%

VC

PVClTPU

Blends

143

Table 5 Mechanical Properties of the ModifierPVC Blends ~

Blends 82.2 Hardness (“ShA) 12.0 (MPa) Tensile strength Ultimate elongation(%) 38 47 deformation(%) Permanent

86.2 13.5 362 53

PVC:Modifier:DOP:Additives= 100:75:20:4

(by weight).

93.4 29.2

PVC modified by TPU has improved abrasive resistance, improved impactresistance (especially at low temperatures), improved oil resistance, improved heat and UV light resistance, and improved adhesive capability with other materials. TPU modified by PVC has improved processability, improvedmoisture resistance, improved flame resistance, improved heat resistance, lower costs, and lower abrasion coefliciency. B.

CommercialPVClTPUPolymericBlends

Several commercial products of PVC/TPU blends are available. The BF Goodrich ChemicalGroup has a PVC/ TPU blend based on their Estane series TPUs. For example, their Estane 54620, a polyester-based TPU with a “ShA85 hardness, shows excellent compatibility with flexible PVC.The blends are produced by mixing PVC, TPU, plasticizer, stabilizer, and lubricant inatwinscrew extruder. These polymeric blends showintermediate mechanical properties between PVC and TPU. Dianippon Ink & Chemical Company (DIC) manufactures the Pandex series of TPUs that are used to make polymericblendswithPVC. These polyblendsshow comparable mechanicalproperties to others. Germany’s Beyer Chemical Company also has similarproducts. The relatedinformation about these commercial products can be obtained from the manufacturers. C.

Others

A few noncommercial PVC/TPU systems have been reported. Makarof et al. [l31 reported a polymer blendof PVC and polyester-based TPU. The TPU series were made from a polyesterglycol (composed of butanediol

1

74.2 22.1 460 30

23 23

and hexanediacid),MDI, and 1,4butanediol as chain extender. The polymeric blends were tested on various mechanical properties and generally showed good prop erties. A series of polyester-based TPU (566TPU series) were synthesized in our lab and usedto blend with PVC to manufacture a modified PVC material for medical uses [14]. Morphological studies showed that 566TPU has very good compatibility with PVC. Detailed mechanical and electronic property tests were also conducted. Some of the data are providedin Tables 3 and 4. These data show that the 566TPU/PVC polymeric blend has good mechanicalproperties, especially at low temperatures. Other tests showed very good oil resistance of this material.Also, the migration rate of plasticizer is only one-fourth of that of commercial medical grade flexible PVC material.

D. The Comparison Between TPU and Other Polymeric Modifiers of PVC A series of tests [l51 were conducted to compare three types of polymeric modifiers for PVC: Du Pont’s Elvaloy 741, a copolymerof ethylene, vinylacetate, and carbon monoxide;Goodyear’sChemigum P83,a copolymer of butadiene andacrylonitrile;and 566TPUfrom our lab, a polyester-basedTPU. Some of the results are provided in Tables 5 and 6. These results showed that TPU, represented here by 566TPU, has the best overall results in modifying PVC amongthose tested. It has good compatibility with PVC, and the resulting polymericblends have good mechanical properties suitable for various processing methods.

Table 6 DSC Results of Modifiers and Their Blends with PVC

P83

p83PVC

E741PVC E74 1

566TPU 566TPUPVC control PVC

’ ~

-23.7 TB (“C) Tm (“C)

50.1 T1.7 ~ =I -17.9 -32.8 11.3 - 24.4 T a = 52.3 57.7 - 47.7

PVC:Modifier:Plasticizer:Additives

= 100:75:20:4 (by weight).

~~

-

-

144

VI. APPLICATIONSANDDEVELOPMENTS Although relatively new, PVC/TPU polymeric blends have already found substantial applications in various fields. This is evident by the numerous patents applied for in this area. Since it possesses good properties of both PVC plastics and polyurethane elastomers, it has been used in those areas where PVC and polyurethane have traditionally played dominant roles. For example, it is a very promising replacementfor flexible PVC used for medical purposes and in the food industry [16,17], because it essentially eliminates the concern regarding plasticizer contamination. It has been usedin combination withthe copolymer of butadiene and acrylonitrile(NBR) to make the abrasion-resistant aprons and rolls used on textile machines [18]. A PVC/TPU/ABS blend serves as a substitute for leather [19]. This could have a tremendous impact on the shoe industry. It has also been found to have an application as a building coating [20,21]. This trend will certainly grow and more applications will be found. This in turn should bring new developments in the material itself. REFERENCES 1. G. Pastuska, Polymer Processing andproperties (G. Astarita, L. Nicolais, eds.) Plenum Press, New York,p. 295 (1984). 2. C. D. Huang and J. Chen, PVC (China), 1: 44 (1992). 3. P. W. Albro, Environ. Health Perspect. 65: 293(1986).

Huang and Chen L. A. Utracki, D. J. Walsh, and R. A. Weiss, Multiphase Polymers: Blends and Ionomers (L. A. Utracki and R. A. Weiss, eds.), ACS, Washington, DC, p. 1 (1989). 5. J. H.Hildebrand andR. L. Scott,Solubility of Non Electrolytes, Reinhold, New York (1950). 6. M. Jiang, Physical Chemistry of Polymer Alloys. Sicuan Educational Press, China(1988). 7. C. D. Huang and H. L. Zheng, Plastics Industry, l : 39

4.

(1991). 8. L. E. Nielsen, Mechanical Properties of Polymer and Composites, Marcel Dekker, Inc., New York (1974). 9. X. Q. Jiang, B. Ha, and C. Z. Yang, Synthetic Rubber Industry (China), 16(4): 220 (1993). 10. The Soviet Union Chemistry Abstract, 7C135 (1981). 11. G. Q.Mei and Y. C. Yan, Liming Chemical Engineering (China), 3: 14 (1995). 12. Z. D. Zhong,B. F. Wang, and Y. Jin, Plastics Processing (China), 4: 19 (1992). 13. A. C. Makarof, et al. Plastics (the Soviet Union), I: 26 (1989). 14. C. D. Huang, H. L.Zheng, J. Chen, J. Zhao, andB. Ding,

unpublished work. 15. C. D. Huang and J. Chen, Po1ym.-Plast. Technol. Eng. 33(5): 615 (1994). 16. Sako Eiji (The Green Cross Corp.) British Patent, No. 1550260 (1979). 17. Sako Eiji (The Green Cross Corp.) Austrian Patent, No. 357013 (1980). 18. C. D. Huang and Z. H. Wu, unpublished work. 19. Luo Lukovsky. Czech Patent, No. CS 246466 (1987). 20. Y. F. Chen and Z. D. Cao, PolyurethaneIndustry (China), I : 22 (1992). 21. Y. F. Chen and W. H. Gu, “PU/PVC Building Coating,”

presented at the 7th Chinese National Polyurethane Meeting, Tai Yuan City, 1994.

11 Mechanical Properties of lonomers and lonomer Blends Masanori Hara and J. A. Sauer Rutgers University, Piscataway,New Jersey

1.

NATURE OF IONOMERS

Ionomers are polymers into which ionic groups have been introduced, usually at a concentration of up to 10-15 mol%, onto some of the monomeric units of the hydrocarbon macromolecular chains, or, in the case of telechelic polymers, at one or both of the chains ends [l-31. Ionomers can be synthesized by copolymerization of one monomerwith another that contains acrylic, methacrylic, sulfonicor other type of acid group, or they may be prepared directly from the homopolymer by chemical modifications. The acid groups, whether carboxylic or sulfonic, are subsequently neutralized, partially or fully, by the addition of appropriate neutralizing agents, such as metal hydroxides. The ionic groups of the ionomer interact with one another and thereby alter the cohesion and the physical and mechanicalproperties. Although the precise microstructure of ionomers is still an open question, it is generally accepted that “multiplets,” consisting ofsmall groupings of ion pairs, and “clusters,” consisting of an ion-richsecond phase containing both multiplets and portions of hydrocarbon chains, are present [4]. Strong evidence for the presence of the ionic cluster “phase” comes fromstudies of ionomers by means of small-angle x-ray scattering ( S A X S ) and small-angleneutron scattering (SANS). Bothof these techniques show the presence of a peak in the scattered intensity at low scattering angles, corresponding to a Bragg spacing of 2-5 nm. This so-called ionic peak is attributed to the presence in the ionomer of a dispersed second phase consisting of ionic clusters. As the ion content of an ionomer increases, the proportion of the cluster phase to the multiplet-containing 145

matrix phase rises. Then, at some critical ion content, or ion content range, the ionic cluster “phase” becomes dominant over the multiplet-containing matrix phase. One way of determining the critical ion content for a given ionomer is to measure the dynamic mechanical properties, i.e., the storage modulus (E’) and the mechanical loss (tan a), as a function of temperature. Measurements of this type have been made on a number of ionomers. The mechanical loss, as a function of temperature, usually shows two peaks, one corresponding to the glass transition temperature of the matrix phase, TgI, and the other, at a higher temperature, to the glass transition of the more tightly bound cluster “phase,” Tg2. This type of plot, for the Na-salt of a sulfonated polystyrene (SPS) ionomer is shown in Fig. 1 [5]. As the ion content rises, the maximum intensity of the matrix peak falls whilethat of the cluster peak increases. For this ionomer,the critical ion content at which the cluster “phase” becomes dominant isabout 6-7 mol%. A comparable valueof the critical ion content has been found in another ionomer, viz. poly(styrene-co-sodium methacrylate). However, for a poly(methy1 methacrylate) (PMMA) ionomer, formed from neutralization of a copolymer of methyl methacrylate and methacrylic acid, the critical ion content appears to be about 12 mol% [6]. One reason for the higher critical ion content for the PMMA-based ionomer is that the PMMA has a higher dielectric constant than polystyrene (PS); hence, at any given ion content, the tendency for multiplets to form clusters is reduced. Thus, a higher ion content is required before the cluster “phase” becomes dominant over the matrix phase. The mechanical properties of ionomers, such as their modulus or stiffness, tensile strength and energy-

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temperature of the parent homopolymer. Anillustration of the profound changesthat can be produced is evident from Fig. 2, which shows how the elastic storage modulus of PMMA-based ionomers depends on temperature [6]. Note the dramatic increase of several decades or more inthe rubbery plateau modulusas the ion content, and the extent of the ionic cluster "phase," rises. There is also a shift of the multiplet-containing matrix Tg to higher temperatures. Both of these results are indications that the ionic groups are acting rather like covalent crosslinks, although ionic bondsin ionomers are generally weaker than covalent bonds. Ionomers, however, have an advantage over covalently crosslinked polymers; i.e., they can be processed by conventional molding techniques provided the ion content is not too high.

Ion Content( Mol YO)

Figure l (tan

versus-ioncontent for the multiplet phase and the cluster phaseof Na-SPS ionomers.

to-fracture, induced deformation modes, and lifetimeto-fracture under alternating loading, depend on many factors. These include type of ionomer, ioncontent, processing conditions and thermaltreatment, percent conversion (neutralization), nature of the counterion, and extent of blending, if any [7].In this chapter, our primary interest is in the mechanical performanceof solid ionomers under ambient temperature conditions and in the influence of various variableson mechanical properties. It is pertinent, however, to first note the very extensive changes in properties that occur at elevated temperatures in the rubbery region above the glass transition

10

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Figure 2 ModulusversustemperatureforNasaltsof PMMA ionomers of various ioncontents.

II. INFLUENCE OF ION CONTENT As the ion content of an ionomer is increased, the proportion of the ionic cluster "phase" that is present rises. This leads to an increase in "entanglement strand density." The significance of strand density to mechanical behavior has been pointedout in the literature [8]. One consequence is that a low-strand density polymer, like polystyrene, deforms in tension only by crazing and is susceptible to early fracture by craze fibril breakdown and subsequent crack propagation through the already crazed material. Incontrast, a high-strand density polymer, like polycarbonate, readily undergoesshear deformation andthe presence of shear leads to enhanced ductility and toughness. For a brittle polymer, such as polystyrene, it is known that the entanglement density can be increased by radiation crosslinking. Thisproduces a change in deformation mode from crazing only to combined crazing and shear, and the presence of shear deformation hinders the growth of crazes. Also, as a result of the crosslinking,and the increase in entanglement density, a higher stress is now required to initiate and propagate crazes. Somewhat similareffects on deformation modes may be anticipated to occur in ionomers as a result of ionic interaction and ionic crosslinking. As an indication of the changes in deformation modes that can be produced in ionomers by increase of ion content, consider poly(styrene-co-sodium methacrylate). In ionomersof low ion content, the only observed deformation mode instrained thin films cast from tetra hydrofuran (THF), a nonpolar solvent, is localizedcrazing. But for ion contents near to or above the critical value of about 6 mol%, both crazing andshear deformation bands have beenobserved. This is demonstrated in the transmission electron microscope (TEM) scan of Fig. 3for an ionomer of 8.2 mol% ioncontent. Somewhat similar deformationpatterns have also been observed in a Na-SPS ionomer havingan ion content of 7.5 mol%. Clearly, in both of these ionomers, the presence of a

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Mechanical Propertiesof Blends

Figure 3 TEM micrograph of a deformed thin film of an 8.2 mol% poly(styrene-co-sodium methacrylate) ionomer

cast from THF.

dominant cluster “phase” and a significantlyhigher strand density has led to the development of shear deformation. The changes observed in strained films are a promising indicationthat beneficial changesin mechanical properties may also be realizedin bulk specimens of ionomers. Relatively few investigationsof the mechanical behavior of solid or glassy ionomers have been reported in the literature [7]. In one study of the tensile properties of Na-SPS ionomers, a modest increase of tensile strength was observed at low ioncontents; and a marked increase, to values about 60% higher than for the PS homopolymer, was found as the ion content approached and exceeded the critical range of about 6 mol% [9]. The energy-to-fracture,or toughness, followed a similar trend with the maximum value beingabout 100% above that of the homopolymer. In general, as the ion content is raised, the modulus or stiffness of the ionomer isincreased, as shown by the data in Fig. 2. While the increase is much greater in the elevated temperature range, where the polymer is acting more likea crosslinked rubber, there is still a significant increase in the glassy modulus below Tg. For example, for the PMMA-based ionomerof Fig. 2, the modulus at 30°C is almost 20% above that of the homopolymer for an ionomer havingan ion content of 12.4 mol%. For the

ionomer with ioncontent of 24.7 mol%, the increase in stiffness is about 45%. Another property of glassy ionomers, which has been investigatedas a function of ioncontent, is its resistance to fracture under applied alternating stress. In tests made on Na-SPS ionomers, the fatigue life was observed to increase with ion content and to reach a maximum valueat about 6 mol%. Thus, such mechanical properties as modulus, tensile strength, energy-to-fracture in simple tension, and resistance to fracture under alternating stress, are all enhanced by introduction of sufficient ionic groups into polystyrene so that microstructure becomesdominated by the ionic cluster “phase.” The clusters provide more effectivecrosslinking than simple multiplets and the resulting increase in ionic entanglementstrand density, as the ion content is raised to the critical value, is then sufficient to induce some of the same beneficial changes in properties that are known to occur as a result of radiation crosslinking of PS. In nonrigid ionomers, such as elastomers in which the T, is situated belowambient temperature, even greater changes can be producedin tensile properties by increase of ion content. As one example, it has been found that in K-salts of a block copolymer, based on butyl acrylate and sulfonatedpolystyrene, both the tensile strength and the toughness show a dramatic increase as the ion content is raised to about 6 mol% [IO]. Also, in Zn-salts of a butyl acrylate/acrylic acid polymer, the tensile strength as a function of the acrylic acid content was observed to rise from a low value of about 3 MPa for the acid copolymerto a maximum value of about 15 MPa for the ionomer having acrylic acid content of 5 wt% [l l]. Other examples of the influence of ion content on mechanical properties of ionomers are cited in a recent review article [7]. 111.

EFFECTS OFPROCESSING CONDITIONS

The mechanical properties of ionomers can be appreciably altered by the manner in which the ionomer is prepared and treated prior to testing. Some of the factors that are influential are: the degree of conversion (neutralization) fromthe acid formto thesalt form, the nature of the thermal treatment or aging, the type of counterion that is introduced, the solvent that is used for preparation of thin films, and the presence and nature of any plasticizers or additives that may be present. In the scope of this chapter, it isnot possible to provide a complete description of the influence of each of these variables onthe wide variety of ionomers that are now commerciallyavailable or producedin the laboratory. Instead, one or more examples of the changes in properties that may be induced by each of the processing variables is presented and discussed.

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A.

Degree of Conversion (Neutralization)

In mostionomers, it is customary to fully convert to the metal salt form; but, in some instances, particularly for ionomers based ona partially crystalline homopolymer, a partial degree of conversion may provide the best mechanical properties. For example, as shown in Fig. 4, a significant increase in modulus occurs with increasing percent conversion for both Na and Ca salts of a poly(ethylene-co-methacrylic acid) ionomer; andinboth cases, at a partial conversion of 30-50%, a maximum value, some 5-6 times higher than that of the acid coa subsequent polymer, isobtained and this is followed by decrease in the property [121. The tensile strength of these ionomers also increases significantly withincreasing conversion but values tendto level off at about 60% conversion. The best combination of properties of polyethylenebased ionomers, such as stiffness, strength, transparency, and toughness, are realized at partial degrees of conversion of about 4040% [131. The initial increase in properties is a result of the presence of ionic interactions, which strengthen and stiffen the polymer. There is, however, some lossof crystallinity as a result of the presence of the ionic groups. When the loss of crystallin-

ity becomesgreat enough to begin to counter the beneficial effects of the ionic crosslinking,a maximum in performance is reached; and this is followed by a decline in properties at higher degrees of neutralization due to the loss of the strengthening effect of the crystalline phase. In someionomers, an excess of degree of neutralization may be helpfulto mechanical properties. For example, in a Zn-salt of a poly(buty1acrylate-co-acrylicacid) ionomer (5.4 mol%), tensile strength was observed to increase steadily with degree of neutralization and to reach a maximum value, some 2.5 times that of the butyl acrylate homopolymer, when there was about 60% excess of neutralizing agent [ll]. As another example, at 100% excess of neutralizing agent in a Na-salt of an SPS ionomer, the tensile strength was increased by about 20%, andthe average lifetime-to-fracture under an alternating stress of 16.5 MPa, was foundto be over 4 times that of a fully neutralized sample[14]. Additional examples are cited in a recent review [7]. Possible reasons for the increase in performance whenan excess of neutralizing agent ispresent are: strengthening of the ionic aggregates by the excess ions, neutralization of any unneutralized acid groups that were initially present, and the presence of the excess neutralizing agent as small, reinforcing second phase particles. B. ThermalTreatment

04 0

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Figure 4 Secant modulus versus percent conversion for Na and Ca salts of an ethylene/methacrylic acid ionomer.

The properties of ionomers are strongly influenced by the manner in which theyare produced and by the thermal treatment that is provided. For a given ionomer to reach its equilibrium microstructure, there must be sufficient time at an elevated temperature for ionic groups to interact and for ionic aggregatesto form. Various examples to illustrate the possible effects of thermal treatment and physical agingare presented. For example, in a compression molded sampleof a Na-SPS ionomer of 2.5 mol% ion content, the storage modulus was found to drop sharply above the matrix glasstransition temperature and no rubbery plateau modulus was achieved. However, when this same ionomer was given an additional heat treatment (24 h at 20O0C), an extended rubbery plateau region was developed aand small mechanical loss peak, indicative of the presence of a cluster “phase,” was observed [15]. Thermal treatment and the nature of the casting solvent can also affect the deformation modes achievedin strained films of ionomers. For example, in films cast from polar dimethylfomamide (DMF), the solvent interacts with ion-richclusters and essentially destroys them, as is evident form absence of a second, higher temperature loss peak in such samples. As a result, even in a cast DMF sample of Na-SPS ionomerof highion content (8.5 mol%), the only deformation modeobserved in tensile straining is crazing.However, when these films are given an additional heat treatment (41 h at 210°C), shear

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Mechanical Properties of Blends

bands are seen in additionto crazes [7]. This is an indication that the ionic clusters have been restored. Additionalconfirmationcomesfromdynamicmechanical measurements, which show the presence of a second elevated temperature loss peak. Studiesof PMMA-basedionomersalso demonstrate the influence of thermal treatment ondeformation modes (16). For Na salts of PMMA-based ionomers of 6 and 12 mol% that were cast from DMF, only crazes were observed on straining. However, after an additional heat treatment (48 h at 160"C), whichalso removes anyDMFsolvent that is present, shear deformation zones are induced. Hence, the ionic cluster "phase," which was destroyed by the polar solvent, has been restored by the heat treatment. For partially crystalline ionomers, such as those based on copolymers of ethylene and methacrylic acid, even time or aging at room temperature can have an effectonmechanical properties. For example, upon aging at 23"C, the modulus of the acid formof the copolymer increased 28%, while in the ionomer form, the increase ranged up to 130%, with the specific gain in modulus depending on the degree of conversion and on the counterion that was present [17].

C. Nature of Counterion The type of counterion present in an ionomer may, or may not, have a significanteffect of properties. For polyethylene-based ionomers, where the presence of crystallinity hasan appreciable effect on properties, the type of counterion present does not appear to have a significant effect on either modulus or tensile strength, as Fig. 4 indicates. However, in amorphous ionomers, the effects of changing the counterion from a monovalent one, as in Na or K, to a divalent one, such as Ca, may be appreciable. As one example, in thin films of Na or K salts of PS-based ionomerscast from a nonpolar solvent, THF, shear deformation is onlypresent when the ion content is near to or above the critical ion content of about 6 mol%; and the TEM scan of Fig. 3, for a sample of 8.2 mol% demonstrates this; but, for a THF-cast sample of a divalent Ca-salt of an SPS ionomer, having only an ion content of 4.1 mol%, both shear deformation zones and crazes are developed upon tensile straining in contrast to only crazing for the monovalent K-salt. This is evident from the TEM scans of Fig. 5. For the Ca-salt, one sees both an unfbrillated shear deformation zone, and, within this zone, a typical fibrillated craze. The Ca-salt also develops a much more extended rubbery plateau region than Na or K salts in storage modulus versus temperature curves; and this is another indication that a stronger and morestable ionic network is present when divalent ions replace monovalent ones. Still another indication that the presence of divalent counterions can enhance mechanical properties comes from

P

Figure 5 TEM micrographs of deformed thin films of an SPS ionomer having an ion content of 4.1 mol% and cast from THF K salt (a), and Ca salt (b).

study of the lifetime-to-fracture under alternating stress conditions. For SPS ionomer samples, with an ion content of 4.1 mol%, divalentCa-salts exhibited an average lifetime-to-fracture that was 2.8 times greater than that achieved by monovalent K or CS salts of the same ion content [NI. The combined effects of a divalent Ca counterion andthermal treatment can be seen from studies of PMMA-based ionomers [16]. In thin filmsof Ca-salts of this ionomer cast from methylene chloride, and having an ion content of only 0.8 mol%, the only observed deformation wasa series of long, localized crazes, similar to those seen in the PMMA homopolymer. When the ionomersamples were subject to an additional heat treatment (8 h at lOO"C), the induced crazes were shorter in length and shear deformation zones were present. This behavior impliesthat the heat treatment enhanced the formation of ionic aggregates andincreased the entanglement strand density. The deformation pattern attained is rather similar to that of Na salts having an ion content of about 6 mol%; hence, substitution of divalent Ca for monovalent Na permits comparable deformation modes, including some shear, to be obtained at much lower ion contents.

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D.

Influence of Plasticizers

In the preparation and processingof ionomers, plasticizers may be addedto reduce viscosity at elevated temperatures and to permit easier processing. These plasticizers have an effect, as well, onthe mechanical properties, both in the rubbery state and in the glassy state; these effects depend on the composition of the ionomer, the polar or nonpolar nature of the plasticizer and on the concentration. Many studies have been carried out on plasticized ionomers and onthe influence of plasticizer on viscoelastic andrelaxation behavior and a review of this subject has been given [19]. However, there is still relatively little information oneffects of plasticizer type and concentration on specific mechanical properties of ionomers in the glassy state or solid state. Two recent studies in this laboratory have examined the effects of a polar and a nonpolar plasticizer on the properties of a PS-based ionomer and on a PMMA-based ionomer each of which was given an ion content in the critical range where the matrix phase and the cluster “phase” had comparable magnitudes. Results of this study on the Na-SPS ionomer (5.5 mol%) are in accordance with the general concept that a nonpolar plasticizer, such as dioctyl phthalate (DOP), selectivelyplasticizes the multiplet-containingmatrix phase, thereby reducing Tgl, but has only a modest effect on the ion-rich cluster “phase.” The polar plasticizer, glycerol, acted in an opposite fashion. It selectively plasticizedthe cluster “phase,” and essentiallydestroyed it, while having only a small effect on the matrix phase. The effects of the same two plasticizers on a NaPMMA based ionomer (12.4 mol%) are somewhat different. For a concentration of 22.4 wt%, the nonpolar DOP reduced the matrix Tgl about 60°C but the cluster Te;! was reducedby only 30°C. Thus, the nonpolar DOPacts in a rather similar fashionfor the PMMA ionomer as for the PS-based one. This is not the case though for the polar plasticizer, glycerol.It, at comparable concentration (19.4 mol%) to that of the nonpolar DOP, acts as a “dual” plasticizer for the PMMA ionomer, decreasing both the Tglof the multiplet-containing matrixphase and the Ta of the ionic cluster “phase” by essentially the same amount, viz. -60°C. Evidently, the higher dielectric constant of PMMA vs that of PS, together with the polar nature of the ester side group in PMMA, causes the polarglycerol to interact more strongly with the backbone chains of PMMA than of PS. The presence of plasticizers not onlyshifts the glass transition temperatures to lower values but it can cause significant changes inthe glassy modulus. Asone example, it has been observed in the Na-PMMA ionomer (12.4 mol%), that the modulus in the solid state (measured at 30°C) decreased at the rate of 2.5% per wt%of DOP but at a higher rate of 7.7% per wt% of glycerol. Thus, for essentially comparable amounts of plasticizer, the polar glycerol has a much greater effect on the properties of

PMMA ionomers in the glassy state than does the nonpolar DOP. The effects of glycerol and DOP, at 10 wt% concentration, on tensileproperties of a glassy Na-SPS ionomer have been examined. Addition of DOP decreased the modulus, the tensile strength, and the toughness. Glycerol at the same concentration, increased the elongation to fracture and the toughness by about 30%, while slightly reducing the stiffness and the tensile strength. The influence of plasticizer concentration on the tensile properties of the same ionomerhave also been explored, over the range from0-30 wt%. With increasingconcentration, the modulusand tensile strength steadily decreased but the strain-to-fractureand the toughness rose and reached maximum values at 15 wt%. At this concentration, the energy-to-fracturerose over 70%. These examples illustrate the wide range in properties that can be realized in ionomers, not onlyby control of ion content and heat treatment, but by selection and control of an appropriate plasticizer. IV. IONOMERBLENDS A.

lonorner/Hornopolyrner

The mechanical properties of blends of ionomers with homopolymers have not beenextensively studied but a discussion has been given of the available information [7]. The purpose of blending is to enhance one or more properties of the homopolymer. It has been discovered that, in several blendsof an ionomer witha suitable polymer, synergistic effects can be realized, i.e., the measured property of the blend becomes significantly higher than values expected on the basis of the simple rule of mixtures. As one example, the influence of blend composition onthe tensile strength of a blend of a sulfonated butyl rubber (SBR) ionomer and polypropylene (PP) is 6 [20]. At all blend compositions the tensile shown in Fig. strength lies above the rule of mixtures value and it reaches a maximum value of 38 MPa at a composition containing about 30 wt% of the SBR ionomer component. The enhanced performance of the blends may result from the presence of two separate interpenetrating networks, one arising from the presence of crystallites and the other from the ionic crosslinking of the ionomer component. Another factor may be good adhesion between twophases resulting from some degree of affhity as a result of common methyl groups in both components. In support of this, in similar blends of SBR with high-density polyethylene (HDPE), which lacks methyl side groups, no synergistic effects are found. Another exampleof favorable synergistic effects in ionomerhomopolymer blends isevident from a study of the tensile properties of blends of an SPS ionomer with PS. Over mostof the composition rangethese two polymers are incompatible. For small additions of the SPS ionomer to PS, TEMstudies of cast thin films showthat

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Mechanical Properties of Blends 0.85-

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Figure 6 Tensile strength versus percent polypropylene in SBR ionomerPP blends.

the ionomeris present as well-bonded,sphericalshaped, small, dispersed particles whose size increases with ion content. Upon straining such films, crazes develop inthe PS matrix and the SPS particles in the crazed regions undergoplastic deformation and becomeappreciably elongated [21]. Also, a higher stress is required to produce craze fibrillation withinthe ionomer particles as a result of the presence of ionic crosslinking. A reinforcing effect of dispersed ionomer particles in similar blends of Na-SPS ionomers and PS is also manifest from measurements of the energy to fracture (toughness) of bulk, compression molded specimens of the blends. This is evident from the toughness vs ionomer content data shown in Fig. 7 for specimens having various ioncontents in the ionomer component [22]. For all three ion contents, there is a marked rise in toughness with increasing ionomercontent with maximumvalues, which rise with increasing ioncontent, being reached at about a 10% addition of ionomer. The measured tensile strengths follow a similar trend. An example of this for a blend in whichthe ionomer component has an ion content of5.26mol% is given in Fig. 8. For both tensile strength and toughness, the maximum values achieved are well above values anticipated on the basis of the rule of mixtures (dotted lines in Figs. 6 and 8). As another example, a synergistic enhancement of tensile strength has been reached in blends of a sulfonated polyacrylonitrileterpolymer (SPAN) witha polyurethane (PU) cationomer [23]. Maximumenhancement wasachieved at a blendcomposition of30170 (PU/ SPAN). At this blend composition, the tensile strength was raised froman initial valueof 78.5 MPato 196 MPa and the strain-to-fracture was at its highest. The en-

Figure 7 Toughnessversusionomercontent(wt%)for Na-SPS ionomer/PS blends.

hancement of mechanical properties is attributed to the strong ionic interactions between ionic groups with opposite sign, which are attached to the respective backbone chains. In a partially crystalline homopolymer,nylon 6, property enhancement has been achieved by blending with a poly(ethy1ene-co-acrylicacid) or its salt form ionomer [24]. Bothadditives proved to be effective impact modifiers for nylon 6. For the blends of the acid copolymer with nylon 6, maximum impact performance was obtained by addition of about 10 wt% of the modifier and the impact strength was further enhanced by increasing the acrylic acid content from 3.5 to 6%. However, blends prepared using the salt form ionomer (Surlyn@:9950-Zn salt) instead of the acid, led to the highest impact strength, with the least reduction in tensile

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strength, at a blend composition containing 25 wt% of the ionomer. B. IonomedSmall Additives

Blending of certain additives to an ionomer can result in both a plasticizing action, thereby enhancing elevated temperature processing of the blend, and also in a reinforcing action, thereby enhancing the strength and stiffness properties of the composite at ambient temperatures. This dual action occurs, for example, when zinc stearate (ZnSt2), whichis a crystalline materialwith a lowmelting temperature, is added to the Zn salt of a sulfonated ethylene-propylene-dieneterpolymer (SEPDM) rubbery type ionomer. Upon melt mixingthe two components, the zinc stearate penetrates and attacks the ionic aggregatesof the SEPDM ionomer. This lowers the melt viscosity andpermits easier processing. In this respect, the zinc stearate additive acts rather as a conventional plasticizer. However, when the stress-strain properties of the blends are examined at ambient temperature, the zinc stearate, which is present in the form of dispersed crystallites, acts as a strong reinforcing agent.This is evident from Fig. 9 where the stress-strain response of a 50/50 blend is comparedto that of the EPDM homopolymer,a sulfur vulcanized EPDM (X-EPDM) and the Zn-SEPDM ionomer [25]. While the ionomer itselfshows significant increases in modulus and strength as compared to the homopolymer or to the vulcanized EPDM, it isevident that the greatest enhancement of strength, stiffness, and toughness (area under the stress-strain curve) occurs for the ZnSt2/Zn-SEPDM blend.The beneficial effects on mechanicalproperties are a result of the reinforcing

action of the ZnSt2 crystalline phase and the network of strong ionic interactions that are present.

V. SUMMARY The mechanical properties of ionomers are generally superior to those of the homopolymer or copolymer from which the ionomer has been synthesized. This is particularly so when the ion content is near to or above the critical value at which the ionic cluster phase becomes dominant over the multiplet-containing matrix phase. The greater strength and stability of such ionomers is a result of efficientionic-type crosslinking and an enhanced entanglement strand density. Aside from ion content, a wide range of properties is availablein ionomers by control of various processing variables, such as degree of conversion (neutralization), type of counterion, plasticizer content and thermaltreatment. Various examples illustrating possible effects of these variables on mechanical relaxation behavior and on such mechanical properties as stiffness, strength, and time- or energy-to-fracture have been given. Blending of ionomers with other homopolymers is also one means of enhancing mechanical performance. Frequently, in ionomer/polymer blends, synergistic effects are realized and properties may be significantly increased over anticipated values based on the rule of mixtures. This area of study has not been extensively explored andthe probability clearly exists that new materials and new blends, having even a greater degree of property enhancement, will become availablein the near future. ACKNOWLEDGMENTS Our research studies on ionomers and on ionomer blends has received financialsupport from the U.S. Army Research Office and ACS-PRF, whoseassistance is gratefully appreciated. Our appreciation is also expressed to former graduate students, Drs. P. Jar, M. A. Bellinger, and X. Ma, who made significantcontributions to some of the research results reported herein. REFERENCES Holliday (ed.), Ionic Polymers, JohnWiley & Sons, New York (1975). 2. A. Eisenberg and M.King, Ion Containing Polymers, Academic Press, New York (1977). 3. L. A. Utracki and R. A. Weiss (eds.) Multiphase Polymers-Blends and Zonomers, ACS Symp. Ser., 395 (1989). 4. A. Eisenberg, B. Hird, and R.B. Moore, Macromol., 23: 1. L.

v

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Figure 9 Stress-straincurves for EPDM,vulcanized EPDM,Zn-SEPDM ionomer,and 50/50 blendofZnSEPDM and ZnStz.

5.

4098 (1990).

M. Hara, P. Jar, and J. A. Sauer, Polymer, (1991).

32: 1622

Mechanical Properties of Blends 6. X. Ma, J. A. Sauer, and M. Hara, Macromol., 28: 3953 (1995). 7. M. Hara and J. A. Sauer, Rev. Macromol. Chem. Phys., C34(3):325 (1994). 8. E. J. Kramer and L. L.Berger, Adv. Polym. Sci.,91/92: l(1990). 9. M.Bellinger, J. A. Sauer, and M. Hara, Macromol., 2 7 1407 (1994). 10. R. D. Allen, I. Volger, and J. E. McGrath, ACS Symp. Ser., 302: 79 (1986). 11. H. Xie and Y. Feng, Polymer, 2 9 1216 (1988). 12. S. Bonotto and E. F. Banner, Macromol., 1: 510 (1%8). 13. R. W. Rees, Polyelectrolytes (K. C. Frisch, D. Klemper and A. V. Paris, eds.), Technomic Publ. Co., Westport, Connecticut, p. 177 (1976). 14. M. Hara, P. Jar, and J. A. Sauer, Macromol., 23: 4964 (1990). 15. M. Hara, P. Jar,and J. A. Sauer, Polymer, 32: 1380 (1991).

153 16. X. Ma, J. A. Sauer and M. Hara, Macromol., 2 8 5526 (1995). 17. E. Hirisawa, Y. Yamamoto, K. Tadano, and S. Yano, Macromol., 22: 2776 (1989). 18. M. Hara, P. Jar, and J. A. Sauer, Macromol., 23: 4465 (1990). 19. C. G. Bazuin, ACS Symp. Ser., 395: 476 (1989). 20. H. Xie, J. Xu, and S. Zhao, Polymer, 32: 95 (1989). 21. M. Hara, M. Bellinger, and J. A. Sauer, Polymer Intern, 26: 137(1991). 22. M. Bellinger, J. A. Sauer, and M. Hara, Macromol., 2 7 6147 (1994). 23. Y. Oh, Y. Lee and B. Kim, J., Macromol. Sci.-Phys., B33: 243 (1994). 24. R. D. Deanin, S. A. Orroth and R. I. Bhagat, Polym. Plast. Technol. Eng., 29: 289 (1990). 25. I. Duvdevani, R. D. Lundberg, C. Wood-Cordova, and G. L.Wilkes, ACS Symp. Ser., 302: 184 (1986).


12 Metallocene-Based Polyolefins: Product Characteristics Virendra Kumar Gupta Indian Petrochemicals Corporation Ltd., Gujarat, India

1.

INTRODUCTION

broader thermoplastic market as more understanding of the differentcommerciallysignificant aspects are realized.

An enormous interest in the synthesis of polyolefins has been created in recent vears due to the discoverv of a new class of catalysts called metallocenes [l-31. These 11- METALLOCENE CATALYST SYWEMS catalysts produce new polymers, such as syndiotactic polypropylene, and cycloolefin copolymers, or can be Metallocene consists of a transition metal atom sandused in the field of already existing polyolefins,such as wiched between ring structures to form a sterically hindlinear low-densitypolyethylene (LLDPE), high-density ered site [6]. The main interest in the use of metallocenes polyethylene, isotactic polypropylene, and ethyleneas catalysts for polyolefins wasstarted in the late 1970s propylene rubber. The most remarkable feature of these when Kaminsky’s group discovered that a small addition emerging catalyst systems is the fact that all metallocene of water in a metallocene-trimethylaluminum system sites produce polymer chains with virtually the same gives highactivity for ethylene polymerization [7]. Since architecture (Fig.1) as comparedwith the different then various types of metallocene-alumoxane systems structures of polymer chains obtained with traditional have been investigated for polyolefin synthesis (Table multiple site heterogeneous catalysts. This characteris2) [l]. These systems show high activities, good copolytic opens up many possibilities for producing tailored merization behavior, and excellent control of stereopolyolefin materials with controlled molecular weights, regularity. However, homogenous systems produce consistent comonomer contents, desired molecular polymers withpoor morphology leadingto their nonutiliweight and chain branching distribution, and control taczation in gas or slurry phase processes. Supported ticity pattern, etc. New moleculararchitecture of homometallocenes are being developed to overcome these geneous polymer molecules coupled with the compatdrawbacks. Inorganic supports, such as silica or magneability of metallocene systems with existing polyolefin sium dichloride, have been used for the incorporation production processes has led to faster commercializaof metallocenes. Another class of catalysts, termed cation of metallocene-based polyolefins (Table 1) [4]. To tionic metallocenes, have also been used [1,8]. Such date, most efforts have been devoted to catalyst develcatalysts provide a hope for development of alumoxaneopment and their evaluation for polyolefin synthesis. free commercial c2talysts. Presently, attention has also been focused on other aspects of the new generation of polyolefin materials, such as processabilityand additive package for various end 111, POLYETHYLENE use applications [5]. Most early applications have been in specialitypolymer markets wherevalue-addedandNovelpolyethylene materials havebeen synthesized by higher priced new polyolefins can compete. The metallo- copolymerization of ethylene with l-butene, l-hexene, cene-basedpolyolefins are expected to compete in a and/or l-octene usingmetallocene catalyst systems [9]. 155

Gupta

156 Table l GlobalCapacityforMetallocene-Based

(a) Metallocene-Single-Site

Polyolefins Catalyst support

Capacity (millions of year) per pound Location Company

(b) Ziegler-Natta-Multi-Site Catalyst support

Med MW/rned branching

LOWMW/high branching

High MW ow branching

Figure 1 Metallocene-single-site versusZiegler-Natta multi-site catalysts. (From Ref. 34. Reprinted with Dermission from Chemical Engineering, McGraw-Hill, Ind., New York, 1993.)

Metallocene and monocyclopentadienyl-amide derivatives of titanium complexes (Fig.2) [2] randomly incorporate ethylene and alpha olefins to give polymers with narrowmolecularweight distribution, highercomonomer contents, and good compositional homogeneity [3].A wide variation in copolymerization efficiency of metallocenes has been observed. Stereorigid metallocenes readily incorporate alpha olefins as compared with nonstereorigid metallocene systems. Dimethyl silyl bridge amido cyclopentadienyl titanium complexes are found to be less restricted in terms of the nature of the comonomer and incorporate a much higher amount of even higher alpha olefins. These homogenous polyethylene polymers are structurally, compositionally, thermally, morphologically, and opticallydistinct materials from traditional heterogenous catalyst-based products. The solution, slurry, and gas phase production processes have been utilizedfor polymer synthesis [9,10]. The so-

Polyethylene Dow Plastics Dow Plastics 125 Exxon Chemical Mitsubishi Nippon Petrochemicals Ube Industries Total Polypropylene BASF 26 Chisso Exxon Chemical Hoechst 220 Mitsui Toatsu Total Polycyclicolefins Dow Plastics Hoechst Pilot Mitsui Petrochemicals

U.S.

Spain U.S. Japan Japan Japan

Germany Japan

us.

Germany Japan

us.

Germany Japan

250 253 220 110

44 1,002

44 220 165

Pilot 7

Ref. 4. (Reprinted with permission from Chem. & Eng. News, Sept 11, 1995, p15; American Chemical Society.)

Source:

lution phaseprocess has mainly been used commercially for the production of different grades of polymers (Table 3). The properties of metallocene-based LLDPE by solution and gasphase processes are shown inTables 4 and 5 PI. LLDPE with narrow molecular weight distribution exhibits a lower, a sharper melting point [3],better hot tack and heat seal properties as well as higher clarity and better impact resistance (Fig. 3), tensile strength (Fig. 4) [ll], and lower levels of alkane-soluble components. The most distinguishingcharacteristic of metallocene-based LLDPEs is that they are not restricted by the current immutable property relationships that are

Table 2 MetalloceneCatalystSystems

Catalyst Nonstereorigid metallocene Stereorigid metallocene Supported metallocene Cationic metallocene Monocyclopentadienyl metallocenes PE = polyethylene; PP = polypropylene.

Cocatalyst Alumoxanes Alumoxanes Trialkyl aluminudAlumoxanes

-

Alumoxanes

PE, atactic PP PE, PP, cyclic polyolefins PE, PP PE, PP PE. PP

157

Metallocene-Based Polyolefins Table 4 Comparison of Properties of LLDPEs

Produced by Metallocene Catalyst and Conventional Ziegler-Natta Catalyst ~

~~~

Property Index

+

7r

MAO

(cm)

+

CpZZrCIz MAO silica

+

Figure 2 Differentmetallocenetypesforpolyethylene. (From Ref. 2.)

~~

~~~

StandardExxpol EX-l01

Melt Density 0.926 flow Spiral 78.1 Tensile strength (MPa) Yield 16.1 7.4Break Elongation (%) 21 Yield Break 77 502 Tensile impact (J/m) Total energy Impact (J) 459 Flex modulus (MPa) 2.2(%) Shrinkage

~~

polyethylene

50

55

0.923 69.9 12.1 7.9 25 275 715 12.1 363 2.0

4.2

Source: Reprinted fromRef. 9, with permission from Elsevier Sci-

commonly applied for Ziegler-Natta polymeric products [2,12]. It includes the control of melt tension of polymer independentof melt flowrate (Fig. 5) and nonexistence of a relationship between melt flow characteristic and molecular weightdistribution. It is possible to synthesize one density grade of metallocene polyethylene with different melt flow rates (Fig. 6) [13]. These developments indicate that metallocene catalysts have provided the opportunity to tailor polymer performance to specific applications-both in terms of processability and end use characteristics. Narrow molecularweight distribution, whichis characteristic of metallocene-based polyethylene (Fig. 7), causes processing difficulty in certain applications due to increased melt pressure, reduced melt strength, and meltfracture [ 14,151. This problem can be overcome by blending the metallocene polymer with other prod-

ence Ltd., Kidlington, U.K.

ucts, such as LDPE, processing aids, etc. Recently, Dow Chemical hasintroduced a new family ofethylenebased branched homogenous polymers using insite technology [2,3,16]. Using this technology, the monocyclopentadienylmetallocene catalyst and process parameters hold the short chain branching level at a precise target while increasing the long chain branching level. This results in the creation of a narrow molecular weight and the homophasic product characteristics of a metallocene catalyst, but with lower melt viscosity and higher melt strength than observed in the homogenous linear

Table 5 Properties of Gas Phase LLDPE (Hexene-l

Comonomer) Using Supported Metallocenes Table 3 Metallocene-BasedPolyethyleneProcesses

Process (company)

MI Comonomer Density range (g110 (glcc)min)

Solution 0.880-0.950 0.5-30 Pow) Solution 1-100 (Exxon) 1 Gas phase (Mobile) 2.3 Gas phase (BP Chem.) Source: Reprinted

Property (%)

0.865-0.940

Octene (0-20%) Propylene,

0.918

Hexene Hexene

0.916

Hexene

fromRef. 9, with permission from Elsevier Science Ltd., Kidlington, U.K.

Density (g/cc) Melt index MFR (g110 min) Dart drop (g) Tear strength Butene, (g/mm) Machine direction Transverse direction Extractables (%) Haze (%)

BP chemicals 0.916 2.3

14.1 18.6 0.8

-

Mobile

Mobil's standard polymer*

0.91 1 18 >800

0.918 1 28 180

14.6 19.3 0.6 5.7

13.8 29.5 3.5 18

* Ziegler-Natta synthesized hexene polymer. Source: Reprinted

from Ref.9, with permission from Elsevier Science Ltd., Kidlington, U.K.

Gupta

158

5

1000

t;

500

.~

d

:!+ 0

-

c

h

EXXPOILLDPEs

\\

"

2 - 3 100-

\Conventional I I I I LLDPEs 35 40 45 50 55 Stiffnessbecant modulus,PSI) ~ ( 1 0 0 0 )

0 -

-5.0

Figure 3 Impact strength of metallocene LLDPE (Ethylene-hexene resin, 1melt index). (From Ref. 11.)

-a. E

+

0 VI

9000-

8000-

Exxpol LLDPEs

\

7000- Conventional LLDPEs

c

6oooo.916

'

o.&o

' 0.A4 '

o.&

Density ( g/cm3 )

Figure properties of meta11ocene LLDPE ylene-hexeneresin,transversedirection).(FromRef.

096

Figure 6 Product profile of LLDPE with respect to density and MFR. (From Ref. 13.)

.-VI

n

0.92

0.88

Density ( g /cc I

e

5

0.64

(Eth-

11.)

ethylene-alpha olefin resins. This new class of ethylenebased polymers couples the advantages of metallocene catalyzed compositional purity with the processing behavior of highlybranching low-densitypolyethylene materials. Polymer design parameters (density, molecular weight, short chain branching, long chain branching, etc.) of polymeric materials can becontrolled independently with Dow's technology. This breaks the existing set of rules regarding process, structure, and property relationship. Metallocene-based polyethylene products are expanding the polyethylene market by taking the polymer into new and existing applications [ 11,17-221. Metallocene-LLDPE (Tables6 and 7)has been targetedfor film and packaging applications while high-density and medium-densitypolyethylene grades are used for injection moulding [23](Table 8 ; applications: thin-walled containers and stadium cups) and rotational moulding (Table 9; applications: agriculture and industrial containers, trash containers, tanks, etc.), respectively. Commercial applications of LLDPE are notably in the blown and

- 1 100

rr

Metallocene LLDPE

-High-MT

rnetallocene LLDPE

c

Z- N LLDPE

0. g' 0.1

v

1

I

10

100

M FR , dg/ml'n Figure 5 HighmelttensionLLDPEsfromsupported metallocenes. (From Ref. 2.)

A, Z.N. LLDPE

High pressure LOPE

, 2.0

3.0

4.0

5.0

6.0

7.0

Log molecularweight

Figure 7 GPC curve of LLDPE obtained withmetallocene catalyst system. (From Ref. 13.)

159

Metallocene-Based Polyolefins Table 6 Metallocene-Based Linear Low Density

Table 9 Metallocene-BasedMedium-Density

Polyethylene

Polyethylene Metallocene

Property

Property 0.915-0.925 0.5-3.0 2.1-2.2 114-115

Density MI distribution weight Molecular point Melting Processability mpact Dart (g) MD tear (g) TD tear (g) % Haze % Gloss

140-800 150-400 350-500 3.8-6.0 100-135

system

Density (glcc) MI (dglmin) Molecular weight distribution Melting point (“C) ESCR, ConditionGood A, F-50 hrs Odor Part impact

Metallocene Catalyst 0.936 3.3 2.1 126 >1000

Low Good

Source: Reprinted from Ref.23, with permissionof Phillips Petroleum Company, U.S.A.

Source: Reprinted from Ref.23, with permissionof Phillips Petroleum Company, U.S.A.

Table 7 Metallocene-Based Very Low Density

Polyethylene Metallocene Property Density MI distribution weight Molecular point Melting (“C) impact Dart (g) MD tear (g) TD tear (g) % Haze

1400

0.908-0.913 1.8-2.4 2.2 100-107 600200-240 380-410 4-1 1

Source: Reprinted from Ref.23, with permission of Phillips Petroleum Company, U.S.A.

cast film use, such as stretch film, as well as can liners and heavyduty sacks. Film applicationsfor metallocene polyethylene are intended to maximize physical, optical, and heat-sealingattributes. Metallocene-LLDPEs have exceptional toughness, low levels of extractables, excellent opticalproperties, and outstandingheat seal and hot tack characteristics. Film manufacturers have incorporated this uniqueset of properties into a variety of filmcatalyst It has provided end users with many making processes. advantages such as: (1) increased packaging speeds due to lower seal initiation temperature, higher hottack, and reduced blocking; (2) reduced package failures due to greater toughness and superior resistance to abuse; (3) improvedpackage aesthetics due to lower haze and higher gloss; and(4) improved packagedproduct quality due to reducedpackage-product interactions, lower odor, and extractables, etc. Mostoften the value of metallocene-based polyethylene is not derived fromone characteristic but from a superior group of properties. IV. STEREOREGULAR POLYPROPYLENE

Table 8 Metallocene-BasedHigh-DensityPolyethylene

Property Density MI (dglmin) distribution weight Molecular point Melting (“C) Spiral flow index (190°C, 1500 psi, dglmin) 10) (psi. modulus Flexural Tensile strength 0 Yield (psi) impact Drop

Metallocene catalyst 0.953 46 2.6 130 4.50 19 4000 2.7

Source: Reprinted from Ref. 23, with permissionof Phillips Petroleum Company, U.S.A.

Propylene polymerizesby metallocene catalysts to give a variety of polymeric structures (Fig. 8 ) [13,24]. Stereostructure and other characteristics of polymers are determined by symmetry of metallocene, cocatalyst nature, andpolymerization parameters. Inaddition to atactic and isotactic polypropylenes, new stereoregular polymers such as syndiotacticpolypropylenes and hemiisotactic polypropyleneshavebeen obtained. Stereoblock isotactic polypropylenecontaining isotactic blocks of alternating orientation and stereoblock polypropylene consistingof isotactic blocks alternating with atactic blocks could also be synthesized [25]. Among these polypropylene structures, isotactic polypropylene and syndiotactic polypropylene have presently gained much commercial significance.Atactic polypropylene is also gaining importance as an additive and polymeric material for blends.

Gupta

160

U Atacticpolypropylene

use in a wide range of product applications. Potential areas of metallocene polypropylene application in the existing polymer market are indicated in Table 10 [28].

A. Isotactic Polypropylene Isotactic PP

%Prx2

* Syndiotactic PP

Many modifications in metallocene structures have been incorporated, as shown in Fig.9, to synthesize isotactic polypropylene with a range of properties includingmolecular weight, isotacticity, mechanical properties, etc.

* Hemiisotactic PP

e

ZrC12

Ph2

Stereobloc isotactic PP

W Stereoblock isotactica t a c t i c PP

Figure 8 Stereostructures of polypropylene and its relationship with metallocene structures. (From Ref. 3, with permissionfromElsevierTrendJournals,Cambridge, U.K.)

Polypropylene can be synthesized by metallocene catalysts with tailored characteristics [26,27] such as: (1) different degree of tacticity; (2) higher comonomer content with uniform distribution; (3) controlled molecular weight, width, and shape of molecular weight distribution; and (4) lower isotactic or atatic product fractions. Presently, only a part of these options are used insynthesizingnewgenerationpolypropyleneresins. These polymeric materials show great potential for their

"l

Figure 9 Metallocene structures for polypropylene synthesis. (From Ref. 2.)

Table 10 Potential Areas of Application for Polypropylene

Molecular polypropylene features of Expected processlproduct effects Potential application NarrowmolecularweightdistributionDifficultextrusionexceptwhereNMWDFiberhonwovens resins are used. Lowmeltstrength,goodmeltdrawabilityFiberhonwovens Low extractables Narrow tacticity distribution Food packaging Heat seal layers in composites Absence of high comonomer fraction ease Narrow composition distribution processesing and comonomer effective in films reducing m.p.; sharp m.p. New copolymers with higher alpha Expect superior ultimate properties Tough films, flexible moulding olefins Source: Reprinted from Ref. 28, with permission of Exxon Chemical Company, U.S.A.

161

Metallocene-Based Polyolefins

[l-31. Earlier discovered metallocenes, such as ethylene bridged indenyl zirconocenes, showed high activity for isotactic polypropylene (i-PP)synthesis as compared to heterogenous catalyst systems. However, the properties comparison revealthat zirconocenes-basedi-PP has less stereoregular, lower melting points, narrow molecular weight distribution, and low-molecular weights (Table 11) [3].Replacement of ethylene bridge with a silylene bridge improvedactivity of these systems to at least ten times that of modern Ziegler-Nattacatalysts. However, molecular weights, crystallinities, and the melting point of isotactic polypropylene are found to be well below those of conventional isotactic polypropylenes. Further modifications in metallocenes have resulted in the formation of isotactic polypropylene at a commercially useful molecular weight and isotacticity [29]. The performance data of somemetallocene structures (Fig. 10) suitable for high-molecular weightisotactic polypropylene are shown in Table 12. Some of the important features of metallocenebased isotactic polypropylene are as follows: 1. Chemical nature andhigh activity of metallocene results in the production of isotactic polypropylene with a very low levelof chlorine ( Secondary amine> Primary amine as, sym-DMEDA EDA TMEDA X H z C H r> + C H z ) d "H(zC- H 2 ) A H r RP = k[CHP]0.34[DMT]0.45[MMA]1.0 TMBDA TMPDA TMEDA The departure of dependence of R,, on the concenDimethylamino > Diethylamino tration of CHP from 0.5 order might be ascribed to inducTEEDA TMEDA tion decompositionof ROOHtype to form ROO.radical,

which has very low activity to initiate monomer polymerization [40], but can combine with the propagation chain radical to form the primary radical termination. For the same reason, the order of concentration of TBH was also lower than 0.5 when the TBH-DMT system was used as the initiator in MMA bulk polymerization. But inthe BPO-DMT initiationsystem as shown in Table

3. ESR Studies and End Group Analysis ESR studies on the initial free radicals were carried out byusing MNP(2-methyl-2-nitrosopropane) or DMPO (5,5-dimethylpyrroline N-oxide) as the spin-trapping agent. The reactions are shown as:

chanism

Reaction Polymerization of Vinyl

233

Table 6 Effects of Monoamines on AAM Polymerization

ture system Initiation Series

amine

APS CHsCHzCH2NHz APSIPA" 1 APSIBA" (CH3CH2CH&NHAPSIDPA (CHsCH2CHzCH&NH APSIDBA

36.9 26.6

of

36.4

1.47.

C H ~ C H ~ C1.24 H~CHZNH~ 2.68 2.78

35.9

APS (CHsCHd3N 2 APSlTEA (CH3CH2CHhN APSRPA (CH3CH2CH2CH2)sN APSlTBA

1.20

104

R,

E,

2.33 3.43 2.89 6.25 6.49

1.oo

62.0

3.01 3.62 3.01 3.61

1.oo

X

.

1.oo 1.20

Series 1: Polymerization in water solution at 45°C. Series 2: Polymerization in mixed solvent HzO: CHsOH = 10:1 at 50°C. mom, [AAM] = 1.00 mom. [APS] = [Amine] = 1.00 X a [PA] = 1.92 x mom; [BA] = 1.61 x moa.

-CH2N(CH3)Rrather than (CH3hN ( C H R) showed that the methyl group isthe preferable group for substitution. Meanwhile,a secondary product was also formed and verified through ESR as *CH2CHzN(CH3)2( N , N dimethylaminoethylene radical) from TMEDA, which was consideredto form fromthe scission of the primary radical as follows:

n-F.

(CH&NCH2CH2"CH2 ( D-MO )

I

CH3

The spin adducts of free radicalsand MNP or DMPO were observed by means of an ESR spectrometer. The data of hyperfine splittingconstants were compiled in Tables 9 and 10 [40-42,44,45]. ESR studies on the initialfree radicals revealedthat the monoalkylamino radical RHN., dialkylamino radical RzN., and aminomethyl radical CH2N< oraminoethylidene radical >N(-CHCH3)were obtained from the corresponding primary, secondary, and cyclic tertiary amine. In case of a tertiary diamine such as TMEDA, formation of

(CH3hNCH26H2 +

CH3-N=~~2

Endgroupanalysis on the charge transfer com[40,41] using tetracyanoethylene plex(CTC)method (TCNE) as a strong electron charge transfer complex acceptor to react with a donor molecule such as aromatic or aliphatic amineforms a CTC or ion radical pairs exhibiting a characteristic absorption in UV or visible range. The CTC method canbe used for the analysis of the amino end group of the polymer formed.

Table 7 Effects of Aliphatic Cyclic Amines on Polymerization

system

Initiation

RP X 104

R,

E,, (kJ/mol)

1.oo

62.5 34.3 40.3 29.9 38.0 54.8

MW of PAAM ~

KPS KpSlMP KPSPD KPS/NMMP KPS/NEMP KPS/NEP [KPS] = [Amine] = 1.0 x

2.52 4.11 3.94 4.55 3.91 3.22

1.63 1.S6 1.81

1S 5

1.28

moVL; [AAM] = 1.0 mol&,, 45OC.

3.37 3.10 2.32 2.94 3.38 0.82

Feng et al.

234

Table 8 Effects of the Diamines on AAM Polymerization

Structure Initiation system

APS CH2CH2NH2 1.99 APSlEDA 36.9 CH3NHCH2CH2NHCH3 APSIsy-DMEDA 1.28 APSIas-DMEDA APSIas-DEEDA APSIas-DMPDA APSRMMDA APSRMEDA APSREEDA APSRMPA APSRTMBDA

RP x 104

(mo1L.s) diamineof

2.14 1.97 2.63 24.9 3.47 6.83 9.101.70 (CH3)2NCH2CH2NH2 4.62 2.46 4.841.12 (C~HS)~NCH~CH ~’JH~ 2.98 5.870.72 (CH3)2NCH2CH2CH2NH2 1.33

of E, PAAM(10-6) (kJImo1)

R, 1.m

62.0 16.2 31.6 19.7

4.100.74 (CH3)2NCH2N(CHs)2 2.08 4.94 9.741.64 (CH3)2NCH2CH2N(CH3)2 ( C ~ H ~ ) ~ N C H ~ C H3.47 Z N ( C Z H ~ )6.83 ~ 0.83 4.07 8.011.91 (CH3)2N(CH2)3N(CH3)2 5.061.68 (CH3)2N(CH2)4N(CH3)22.57

APS:ammoniumpersulfate;[APS]

=

[Amine]

=

5.0

X

MW

35.6 22.0 24.2 25.0 27.5

mom, [AAM] = 1.0 m o a , 45°C.

4. Initiation Mechanism of Persulfate-Amine N-methylmorpholine [46], which involvesthe formation Systems of aminium radical as the intermediate step and deproBased on theresults discussed previously, the following tonation to methylradical as one of theactive iniinitiationmechanism was proposed for a tertiary amine,tiators:

Table 9 Hyperfine Splitting Constants of Spin Adduct Formed APSIAmineIMNPor DMPO System

System

Radical trapped by MNP or DMPO

Hyperfine splitting constant (0.1 mT) aaN

aa*

aaN

PAIAPSIMNP BNAPSiMNP EDAIAPSIMNP PDNAPSIMNP EDNAPSIDMPO DPNAPSDMPO

15.46 15.47 15.47 15.47 13.68 13.65

0.67(2H) 1.33 0.67(2H) 1.33 0.67(2H) 1.33 0.67(2H) 1.33 9.87 10.0

1.33 1.33 1.33 1.33 1.28 1.19

MPIAPSIMNP

18.72

0.88(2H)

0.88

PDIAPSNNP

18.61

0.93(2H)

0.93

MPIAPSIDMPO

15.0

9.33

2.13

PyDIAPSIDMPO

15.25

17.6

2.29

NMMPIAPSIMNP

14.16

2.38(2H)

0.80

NEMPIAPSNNP

15.28

NMPIAPSIMNP

15.07

3.20 7.55(2H)

3.20 3.52

aYH

chanism

Reaction POlymeriZatiOn of Vinyl

235

Table 10 Hyperfine Splitting Constants of Spin Adducts Obtained from APS/Amine/MNP Systems

Hyperfine splitting constant (0.1 rnT)

ped Radical System

ihN

aeH

aeN

TMEDA

15.73

9.84(2H)

2.08

TMPDA

15.87 15.0

10.35(2H) 8.0(2H)

3.63

TMBDA

15.0

8.0(2H)

TEEDA

15.57

5.63

as-DMEDA

15.73 15.87 15.60

19.5 10.67 9.73(2H)

2.13

as-DMPDA

16.0 15.36

10.8(2H) 9.47(2H)

2.53

DETA TETA TEPA

14.93 14.96 14.93

aYH

0.80(2H)

0.53

0.53(2H)

2.13 2.16 2.16

In case of secondary amine or primaryamine it would form amino radical as follows:

1

The ethylenediamine derivative [31] possesses higher promoting activities than other diamines. This phenomenon maybe ascribed to thecopromoting effect of the two amino groupson the decomposition of persulfate through a CCT (contact charge transfer complex) formation. So we proposedthe initiation mechanism via CCT as the intimate ion pair and deprotonation via CTS (cyclic transition state) as follows:

236

Feng et al.

tors were studied in great detail in the 1960s and 1970s. The most notable work in this area was the photopolymerization of VCZ, as studied by Shirota et al. [62-641 4 and the photopolymerization of m-methyl-styrene (mMSt), as studied by Irie and Hayashi [65]. VCZ as a strong electron-donating monomer can CTS form CTC with variousstrong electron acceptors in the ground state, such as chloranil (CA), bromanil (BA), tetracyanoethylene (TCNE), maleic anhydride (MAn), tetranitromethane (TNM), and trinitrobenzene (TNB), and can be polymerized in the presence of such acceptors. With many weak electron acceptors, however, usually + no polymerization occurs at room temperature in the SOT dark. Polymerization of these systems are induced or 0J accelerated by photoirradiation since the charge-transfer interaction is greatlyenhanced under irradiation and Moreover, the initiation mechanism of APS/ethylenethe excited charge-transfer complex is rather polar and disecondary-amine and APS/ethylene-diprimary-amine readily gives rise to ion radicals. systems has also been proposed [47]. Studies of the effect of the wavelength of the incident light on polymerization have shown that two possi111. PHOTO-INDUCEDCHARGETRANSFER ble processes produce the active initiating species: (1) POLYMERIZATIONi)F VINYL a stable ground state CTC is excited, and (2) either a MONOMERS donor or an acceptor is excited, followedby charge transfer and electron transfer with a ground state accepThe charge transfer theory was established by Mulliken tor or donor. While the photopolymerization of VCZ[48-531 in the 1950s. Since then a great number of orthiobenzophenone in benzene or toluene is initiated ganic reactions have been foundto proceed throughforthrough local excitation of thiobenzophenone [66], the mation of charge transfer complexes (CTC) [54-571. In polymerization of VCZ-p-quinoid compounds in benthe 1960s, this theory was appliedto the polymerization zene [67] and VCZ-trinitrofluorenone in nitrobenzene field and successfully explained the mechanism of many [68] has been shown to be initiated by selective excitapolymerizations. Someof the earliest studies were pubtion of the ground state CTC. Obviously, only the seclished byScott et al. [58] and Ellinger[59] independently ond process is operative for the system where no chargein 1963.They reported the polymerization of N-vinylcartransfer interaction exists in the ground state. For the bazole (VCZ) as an electron-donating monomer in the VCZ-CAandVCZ-BA systems, the above two propresence of electron acceptors through the formation of cesses are operative, depending on the wavelengths of CTC. Beforethat, Nomsh et al[60,611 had reported the the incident light used [67,69]. polymerization of styrene photoinitiated by anthracene Charge-transfer and electron-transfer processes in and proposedthe formation of anexcited triplet complex photopolymerization systems have been demonstrated as an intermediate, butdidnotmention whether the by the measurement of fluorescence spectra andby complex was-of the charge-transfer type or not. Charge-transfer polymerization has been developed means of flash photolysis.Typically, VCZ-dimethyl terephthalate (DMPT) [70], VCZ-FN, and VCZ-diethyl fuwithin the last 30 years on the basis of the interaction marate (DEF) [71] systems have been shown to form between an electron donor (D) and an electron acceptor exciplexes by local excitation of VCZ.Dynamic quench(A) involved in the initiating and/or propagating proing of the VCZ fluorescence occurred by addingaccepcesses. Such polymerization has attracted a great deal tors. Broad and structureless exciplexes were observed of attention because it isprocessed through a novel type at longer wavelengths in nonpolar solvents. Although no of initiating or propagating mechanism with a charge exciplex fluorescence was observed in the polar soltransfer interaction. This exists throughout in a wide vavents, the quenching of the VCZ fluorescence occurred, riety of organic compoundsas donors or acceptors and indicating the occurrence of charge transfer or electron also is found with low energy due to the charge transfer transfer in the excited singlet state of VCZ [72,73]. With interaction in the formation of the reactive centers. flashphotolysisand laser photolysis techniques, the Therefore, the polymerizationwould be camed out photochemical formation of the transient VCZ cation under more moderate conditions. radical andthe electron-acceptor (A) anion radical (A--) [74-771. The transient absorption has been confirmed A. Vinyl Monomers as Electron Donors spectra obtained for the VCZ-A system in various solCharge-transferphotopolymerizations of electron-dovents consists of two band systems: one is due to A.nating monomers initiated by electron-accepting initiaand the other is due to the VCZ cation radical (VCZ-+).

237

Reaction Mechanismof Vinyl Polymerization

These results indicate that a complete electron transfer is brought about from the charge-transferinteraction and can be summarized [62]as follows: AS + VCZ?

VCZ + A

47: C6H5N02

AS + VCZ?

cationicpolymerization Radical polymerization

HPMA: Hexamethylphosphoramide

B. Vinyl Monomers as Electron Acceptors 1. Aromatic Tertiary Amine and ElectronAccepting Monomer In 1982 Wei et [78,79]studied the quenching of N,Ndimethyltoluidine(DMT) fluorescence by adding the electron-accepting monomerMA or MMA and successfully observed broad and structureless exciplex fluorescences at longer wavelengths in nonpolar solvents for the first time. It has been shown that the excited charge transfer complex canbe formed through two different routes depending on the wavelengths of the light used [80-831. Both routes can bring about polymerization. For example, AN can form CTC with DMT in the ground state, which was verified through UV spectra. This CTC will be photoactivated with 365-nm light to form an exciplex. If 313-nm light is used, the activated DMT will form an exciplex directly withthe ground state AN molecule as follows:

al.

DMT

hv,313nm

L

DMT

Route (1) is referred to as “local excitation” and route (2) as “CTC excitation.” It has been observed that the different routes bring about the polymerization of AN with different kinetic behaviors. A 365-nm light willirradiate the CTC only, and in this case the rate of polymerization for different aromatic tertiary amines descends in the following order: P - C H ~ C ~ H ~ N ( C H ~~) ~- C H ~ C ~ & N ( C H ~ C H Z O H ) Z > P - H O C H ~ C ~ H ~ N ( C HCaHsN(CH3)z ~)~ It is in the same order as the equilibrium constants of CTC ofamine-FN. That is, the stronger the ability of an amine to form CTC with electron acceptors, the faster the rate of photopolymerization. However, under 313nm irradiation, local excitation plays a principal role and the rate of polymerization is observed to descend in a different order [801: P-CH~C~H~N(CH~)~ > C&N(CH3)2 > ~ - H O C H Z C ~ H ~ N ( C H ~ ) ~ >P-CH~C~&N(CH~CH~OH)~ This order agrees with that of quenching constant (kq7) values of the fluorescence of amines by AN. That is, the easier the reaction between an excited aromatic tertiary amine andthe ground state AN, the faster the initiation. Local excitation was also studied for primary and secondary amines under irradiation at 313 nm. The results are summarized in Table 1l . In order to estimate the photoinitiating efficiency of the amines, the measurement was performed at a chosen constant absorbance (0.40) of the reaction mixture. The rates of polymerization were found to be in the following order: CsHsN(CH3)2 CsHsNH2 > CsHsNHCH3 > P - C H ~ C ~ H ~>Np-CH3OCsH4NH2 H~ > m-CH3CsH4NH2

*

I AN

which isconsistent with the order of the quenching rate constant kq values calculated with the Stern-Volmer equation:

Table 11 The Photopolymerization of AN Initiated by Aromatic Amines (at 313 nm)

Amine

E313a

2.2 CaHsN(CH& CaHsNH2 5.2 2.3 GHsNHCH3 P - C H ~ C 3.0 ~H~NH~ 1.6 p-CH3OC6HdNH2 3.8 m-CHGH4NH2

1800 860 1760 1240 2300 1030

Cb x 104

46.0 48.0 48.0 27.3

R,, (mol/L.s) x 10s

3.4 2.6 2.3 2.1 1.8 1.7

7(Wd

2.00 1.93 44.5 1.78 1.78 44.0 1.66 1.09

2.3 2.3 2.7 2.5 2.9 2.5

The molar extinction coefficientof amine at 313 nm measured in DMF-AN solution ([AN] = 3.8 m o a ) . The amine concentration (rnoVL) used in the polymerization at which the absorbanceof the system was 0.40. c Calculated with the Stern-Volmer equation, Ion = 1 + /cq TIAN]in alcohol ([amine] = 2 X IO4 mol/L). Measuredinalcohol([amine] = 2 X lo4 moVL).

kq x

lo-*

Feng et al.

238

Z,/Z = kq7[AN]

+1

This result reveals that exciplex formationplays a principal role in the initiation of polymerization. Since the absorption bandisbroadenedtowardlongerwavelengths as the result of formation of CTC between AN and aniline, a certain concentration of aniline can be chosen so that 365-nm lightis absorbed only bythe CTC but not by the aniline molecule. Therefore, in this case the photopolymerization may be ascribed to the CTC excitation selected. For example, a 5 x mol/Laniline solution in AN could absorb light of 365 nm, while solutions in DMF or cyclohexane with the same concentration will show no absorption. Obviously, in this case the polymerization of AN is caused by CTC excitation. The rates of polymerization for different amines were found to be in the following order (Table 12): The process for initiating radicalformation in aromatic amine-vinylmonomer systems have beenstudied by Feng et al. [80-861 who proposed the formation of an aminium radical as the active state of an exciplex as intimate ion-pair and then a cyclic transition state which then would undergoa proton transfer process of deprotonation leading to the formation of active radical species for initiation as follows:

oi!< CH2

C H C~~ H - C N ( CTC )

Intimate ion pair

Cyclic transition state

1

n+ trender

I 2 . Aromatic Secondary or Primary Amine and Electron-Accepting Monomer Li et al. [87,88] found that aniline willprocess the photopolymerization of ANeither in N,N-dimethylformamide (DMF) solutionor in bulk witha fair rate of polymerization only next to DMT. From UV spectra it is proved that aniline will form a CTC with AN. Using 313-nm radiation that CTC is excitedto an exciplex and polymerization proceeds. N-methylaniline will polymerize AN similarly. The following mechanism was proposed:

U

Exciplex,intimate ion pair

l U

(Amino radical )

H+ transfer ( nonpolar solvent)

chanism

Reaction Polymerization of Vinyl

239

Table 12 The Bulk Polymerization of AN Initiated by Aromatic Amines

Absorbance Amine

in DMF

AN Absorbance in 0.03 0.09 0.21

0

co.01 co.01 ([amine]

=

5

X

RP (m0VL-s) 1.7 2.1 2.8

X X

X

10-4 10-4 10-4

lo-* mol at 365 nm)

R

l

o b ...

H+ transfer

+

=CH,CH,CN

( Amino radical)

Thus, deprotonation of the aminium radical froma secondary or primary aminewill at last form an amino radical instead of an aminoalkyl radical anda -CH.LCH*CN radical. This amino radicalwill then serve as one of the active species for the initiation of polymerization. C. Aromatic Secondary or Primary Amine and Benzophenone

The well-known photopolymerizationof acrylic monomers usually involvesa charge transfer system with carbonyl compound as an acceptor and aliphatic tertiary amine, triethylamine (TEA), as a donor. Instead of tertiary amine such as TEA or DMT, Li et al. [89] investigated the photopolymerization of AN in the presence of benzophenone (BP) and aniline (A)or N-methylaniline (NMA) and found that the BP-A or BP-NMA system will give a higher rate of polymerization than that of the well-known system BP-TEA. Still, we know that secondary aromatic amine would bedeprotonated of the Hatom mostly onthe N-atom so we proposed the mechanism as follows:

Pl

Exciplex

f I

/ \ H

Q,

Q CA

-

'Q OH

Thus, an aminium radical from primary or secondary amine will at last form an amino radical instead of an aminomethyl radical.This amino radicalwill then serve as the only active radical species to initiate the vinyl polymerization. The end group of the polymers, photoinitiated with aromatic amine with or without the presence of carbonyl compound BP, has been detected with absorption spectrophotometry and fluororescence spectrophotometry [W]. The spectra showed the presence of tertiary amino endgroup in the polymersinitiatedwith secondary amine such as NMA and the presence of secondary amino end group in the polymers initiated with primary amine such as aniline. These results show that the amino radicals, formed throughthe deprotonation of the aminium radical in the active state of the exciplex from the primary or secondary aromatic amine molecule, are responsible for the initiation of the polymerization. . The proton transfer mechanism described previously was confirmed somewhat bythe influence of solvent polarity on polymerization. The rate of photopo-

Feng et al.

240

lymerization of ANinitiated by aniline an inacetonitrilecyclohexane mixturehas been studied, and it was found that with an increase in the proportion of acetonitrile, the rate of polymerization first increases because the moderate polar solvent is favorable for exciplex formation, andthen the rate falls as thepolarity increases further. This is ascribed to the dissociation of the exciplex into solvated ion-radicalsin a highly polar solvent. This result shows that the intimate radical-ion pair state of the exciplex, but not the solvated ion-radical state, is responsible for proton transfer:

Exdplex Intimate radlcal pair

Dissociation (Polar solvent)

/

H+ transfer

I

( nonpolar sohrent)

(oit)%b

H

- .. +

+

The polymers initiatedby BP amines were found to contain about one amino end group per molecular chain. It is reasonable to consider that the combination of BP and such polymers will initiate further polymerization of vinyl monomers.We investigated the photopolymerization of MMA with BP-PMMA bearingan anilino end group as the initiation system and found an increase of the molecular weight from GPC and viscometrical measurement [91]. This system can also initiate the photopolymerization ofAN to form a block copolymer, which was characterized by GPC, elemental analysis, and IR spectra. The mechanism proposed is as follows:

.CH,CH,CN

Initiation

No initiation

We have prepared a copolymer-bearing aminoside group and used it either alone or in combination with BP to initiate the photopolymerization ofMMA [89]. The gelpermeationchromatography(GPC)plot of PMMA initiated by the former system showed a bimodal distribution of molecular weightbecause both the radicals produced initiate polymerization as follows:

Thus, BP-aniline may serve as the photoinitiator in the establishmentof some new methods of block copolymerization. ACKNOWLEDGMENT

I

CH2 CH*-N

N

I ! !

U

+

(1)

Initiation

C&-

F"

+ CH~-C.

I

COOCH~

(U)

1

Initiation

The high-molecularweight was assigned tothe PMMA grafted to the copolymer chains and the lowmolecular weight to the PMMA initiated by the MMA radical (11). However, only one molecular weightdistribution peak was observed for the PMMA initiated by the latter system, i.e., in combination with BP, which implies that only aminomethyl radicals are capable of initiating the polymerization.

The authors are indebted to the National Natural Science Foundation of China for financial support of this work and also grateful to those coworkers whose research works have been cited. REFERENCES 1. C. Walling, Free Radicals in Solution, John Wiley & Sons, Inc., New York, (1957), p. 590. 2. G. M. Brauer and H.Argentar, Initiation of Polymerization, F. E. Jr. Bailey, ed., American Chenical Society, Washington D.C., p. 359 (1983). 3. Q. R. Wu, D. F. Yen, and X. D. Feng (S. T. Voong), Preprints, The 4th China-Japan Symposium on Radical Polymerization, October 1-3, 1986, Chengdu, China, p. 156. 4. X. D. Feng, Chinese J . Polym. Sci., 3: 109 (1986). 5. X. D. Feng, Makromol. Chem., Macromol. Symp., 63: 1 (1992). 6. W.X. Cao, J. F. Chung, and S. T.Voong, Polym. Commun., 2: 80 (1978). (in Chinese).

chanism

Reaction Polymerization of Vinyl G. E. Serniuk, U.S. Patent, 2,529,315, Nov. 7, 1950. T. Tanaka, Scientific American, 244(1): 110 (1981). E. Boschetti, U.S. Patent, 4,189,370, Feb. 19. 1980. M. K. Gupta and R. Bansil, Polym. Preprints, 22(2):375

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on the Synthesis and Materials Scienceof Polymers, October 21-24, 1984, Beijing, China, p. 13. K. Y. Qiu, J.Y. Zhang, Y. 3. Hu, andX.D. Feng, Polym. Commun., (Beijing), I: 76 (1985). K. Y. Qiu and T. Zhao, Acta Polym. Sin., I: 72 (1988). J. Fu, X. Q. Guo, K. Y. Qiu, and X. D. Feng, Preprints, Symposium on Polymer Syntheses, Polymerization reaction and Mechanism, Nov. 7-11, 1988, Nanjing, China, p. 7; Acta Polym. Sin., I: 67 (1990). K. Y. Qiu, D. J. Guo, andX. D. Feng, Acta Polym. Sin., 1: 84, (1991). K. Y. Qiu, Z. H. Zhang, and X. D. Feng, Acta Polym. Sin., 2:-178 (1993). K. Y. Qiu, Y. B.Shen, andX.D. Feng, ChineseJ. Polym. Sci., 4: 92 (1986). Q. R. Wu, Y. L. Li, and X.D. Feng, Preprints, Kunming International SymposiumonPolymericBiomaterials, May 3-7, 1988, Kunming, China, p. 127. F. M. Li, L. Wang, and X. D. Feng, Polym. Commun., (Beijing), 2: 144 (1987). F. M. Li, L. Wang, Y. W. Ding, and X.D. Feng, Polym. Commun., (Beijing), 4: 316. F. M. Li, L. Wang, Y. N. Ding, and X.D. Feng, Polym. Commun., (Beijing), 5: 393 (1985). F. M. Li, Preprints, The 1st Japan-China Symposiumon Radical Polymerization, Nov. 5-6, 1980, Osaka, Japan, p. 45. F. M. Li, Z. W. Gu, W. P. Ye, and X. D. Feng, Polym. Commun., (Beijing), 2: 122 (1983). F. M. Li, L. Wang, and X. D. Feng, Polym. Commun., (Beijing), 3: 237 (1984). Z. Y. Zhang, H. R. Zhou, A. R. Wnag, and X.D. Feng, Polym. Commun., (Beijing), 2: 149 (1987). F. M. Li, W. P. Ye, and X. D. Feng, Polym. Commun., (Beijing), 5: 397 (1982). K. Y. Qiu and J. Fu, Chinese J. Polym. Sci., 7(3): 258

(1989). Qiu, andX. D. Feng, Chem. J.Chinese 31. J. Y. Zhang, K. Y. Univ., l(2): 114 (1985). 32. K. Y. Qiu, D. J. Guo, andX.D. Feng, ChineseJ. Polym. Sci., 8: 363, (1990). 33. K. Y. Qiu, J. Fu, X. Q. Guo, and X. D. Feng, Chinese J . Polym. Sci., 8: 188 (1990). 34. K.Y.Qiu, X. Q. Guo, X. L. Wang, and X. D. Feng, Chem. J. Chinese Univ., I(1): 141 (1985). 35. Y. H. Sun, K. Y. Qiu, and X. D. Feng, Kexue Tongbao (Science Bulletin), 28(4): 475 (1983). 36. S . T. Voong and K. Y. Qiu, Kexue Tongbao (Science Bulletin), 17(4):247, (1966). 37. Y. H. Sun, K. Y. Qiu, and X. D. Feng, Scientia Sinica, B27(4): 349 (1984). 38. Y. H. Sun, DoctorThesis, PolymerDivision,Department of Chemistry, Peking University (1984). 74.

241 39. J.Pavlinecand M. Lazar, Collect Czech Chem. Commun.. 42: 5023 (19771. 40. X. Q: GUO,K. Q ~ uand , X. D. Feng, Scientia Sinica, B30f9):897 (1987). 41. X. Guo, docto; Thesis, Polymer Division, Department of Chemistry, Peking University (1987). 42. X. D. Feng andK. Y. Qiu, Advances in Science of China, Chemistry, 4: 39 (1992). 43. K. Y. Qiu, X. Q. Guo, and X. D. Feng, Kexue Tongbao (Science Bulletin), 31(11): 736 (1986).

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Q.

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X. Q. Guo, K. Y. Qiu, and X. D. Feng, Preprints, The 5th Japan-China Symposiumon Radical Polymerization, August 8-10, 1988, Osaka, Japan, p. 33. R. S. Mulliken, J. Am. Chem. Soc., 72:600,4493 (1950). R. S. Mulliken, J. Am. Chem. Soc., 74:811 (1952). R. S. Mulliken, J. Chem. Phys., 1 9 514 (1951). R. S. Mulliken, J. Chem. Phys., 23: 397 (1955). R. S. Mulliken, Rec. Trav. Chim., 75: 845 (1956). R. S. Mulliken, J. Phys. Chem., 5 6 801 (1952). E. M. Kosower, Prog. Phys. Org. Chem., 3: 81 (1%5). R. Foster (ed.), OrganicCharge Transfer Complexes, Eleck Science, London, 1974. R. Foster (ed.), Molecular Complexes,Vol. 2, Eleck Science, London, 1974. S. G. Cohen, A. Parola, and G. Parson, Chem. Rev., 73: 141(1973).

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kawa, andH.Tsubomura, 24th Annu. Meet. Jpn. Chem. Soc., p. 11232 (1971). Y. Shirota, K. Kawai, N. Yamamoto, K. Tada, H. Mikawa, and H.Tsubomura, Chem. Lett., p. 145 (1972). Y. Shirota, K.Kawai, N. Yamamoto, T. Shida, H. Mikawa, and H.Tsubomura, Bull. Chem. Soc. Jpn., 45: 2693 (1972). Y. Shirota, T. Tomikawa, T. Nogami, N. Yamamoto, H. Tsubomura, and H.Mikawa, Bull. Chem. Soc. Jpn., 47: 2099 (1974). Y.Wei, W. X. Cao,and X. D. Feng, KexueTongbao (Science Bulletin), 28: 1068 (1983). Y. Wei, W. X. Cao, and X. D. Feng, Chem. Abstr., 98: 143870t (1983). T. Li, W. X. Cao,and X. D.Feng, Polym. Commun. (Beijing), (2): 127 (1983). T. Li, W. X. Cao, and X. D. Feng, Chem. Abstr., 99: 140446U (1983). T.Li, W. X. Cao,,and X. D. Feng, Polym.Commun. (Beijing), 4: 260 (1983).

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17 Photoinitiation of Free Radical Polymerization by Organometallic Compounds Salah M. Aliwi Mustansiriya University, Baghdad, Iraq

1.

INTRODUCTION

Photopolymerization, in general, can be defined as the process whereby light is used to induce the conversion of monomer moleculesto a polymer chain. One can distinguish between true photopolymerization and photoinitiation of Dolvmerization mocesses. In the former. each chain Gopagation step involves a photochemici process [ 1,2] (i.e., photochemicalchain lengthening process in whichthe absorption of light is indispensable for

2. Through direct excitation of a monomeric o r polymeric molecule or of a molecular complex (A) followed by a reaction producing an initiating species:

each propagation step); whereaslatter, in the only the initial step results from the interaction of light with a photosensitive compound or system [3-61. This chapter deals only with the photoinitiation of polymerization process or photochemical postpolymerization.This may occur in two different ways: 1. Through energytransfer of an excited sensitizer molecule (S) to either a monomer (M) or foreign molecule (A) resulting in the formation of species capable of initiation (e.g., radical):

A

hv

.3

The rate of formationof radical fromthe photoinitiator (the Only khtmabsorbingmolecule in the system) VR generally is given by Eq. (4): O(

IJI-e

&In 1 [In

1)

R. (~ree-rad~as)

(2)

oxctted

The radicals created in (1) and (2) interact withmately,bytermination monomermolecules to produce macroradicals, andulti-chain is produced:

VR.=

*

of these radicals, a polymeric

are where I, is the incident light intensity, [In] and the photoinitiator concentration and its molar extinction coefficient, respectively, and l is the cell path length,

(4)

while (Y is the overall quantum yield of the initiating ;pel cies (K). 243

Aliwi

244

Not all initiating radicals (R') succeed in initiating polymerization, recombination of these radicals in the solvent can decrease the efficiency to a value lower than 1. Detailed kinetictreatment of photoinitiation processes are discussed by Oster and Yang [3]. Photoinitiation of polymerization has played an important role in the early developmentsof polymer chemistry. The main features of this type of initiation are:

(n

1. The rate of formation of the initiating species

and, accordingly, the molecular weightdistribution can be controlled by variation of the incident light intensity. 2. Photoinitiation isan excellent method for studying the pre- and posteffects of free radical polymerization, and from the ratio of the specific rate constant (k,) in non-steady-state conditions, together with steady-state kinetics, the absolute values of propagation (kp)and termination (kt) rate constants for radical polymerization can be obtained. 3. Photoinitiators provide a convenient route for synthesizing vinyl polymers with a variety of different reactive end groups. Under suitable conditions, and inthe presence of a vinyl monomer, a block AB or ABA copolymer can be produced which wouldotherwise be difficultor impossible to produce by another polymerization method. Moreover, synthesis of block copolymers by this route is much more versatile than those based on anionic polymerization, sincea wider rangeof a monomers can be incorporated into the blocks. 4. Photoinitiation witha high quantum yield of radical productionin the visible light is of practical importance for photocuring processes [5,6].

If (P) is terminated by a chain transfer to a solvent or a monomer, a graft copolymer is formed, or, if the termination is from a combination, a crosslinked network polymer isformed. If the pre-existing polymer (B) contains an end group that itself is photosensitive (or can produce a radical by interacting with photoinitiator) and in the presence of a vinyl monomer (A), block copolymer of type AB can be produced if the photosensitive group is on one end of the polymeric chain. Type ABA block copolymer can be produced if the polymer chain (B) contains a photosensitive group on bothends. Free radicals capable of initiating polymerizationof

Oneof the first methods of polymerizingvinyl monomers was to expose the monomer to sunlight. In 1845, Blyth and Hoffman [7] obtained by this means a clear glassy polymeric product from styrene. Berthelot and Gaudechon[8] were the first to polymerize ethylene to a solid form and they used ultraviolet (UV) light for this purpose. The first demonstration of the chain reaction nature of photoinitiation of vinyl polymerization was done by Ostromislenski in 1912 [9].He showed that the amount of poly(viny1bromide) produced was considerably inexcess of that produced for an ordinary chemical reaction. In morerecent years, photoinitiationof polymerization provedto be of immense value in the understanding of the precise nature of polymerization. Several systems used for the initiation of radical polymerizationwere reviewed byOster and Yang [3], Rabek [ 101, and Davidson 15$1. Two types of photoinitiation processes may be considered: A chain reaction polymerization of vinyl monomer, which is usually carried out by a photoinitiator to produce a primary radical ( R ) , which can interact with a monomer molecule(M) in a propagating process to form a polymer chain composed of a large number of monomer units (see Eq. [2] and reaction Scheme [3]. The photografting and photocrosslinking processes in which the photoinitiation iscarried out by pre-existing polymer in the presence of vinyl monomer. The pre-existing polymer may itself be photoactive or, with the presence of photoinitiator, can interact with a polymer to produce a radical onthe backbone of the polymer chain.

vinyl monomerscan be produced photochemically from a wide varietyof substances, such as organic, inorganic, and organometallic compounds one in or two component initiating systems. The photoinitiation of vinyl polymerization by organic compounds (carbonyl, azo, peroxide, disulphide compounds, etc.) orinorganic salts (e.g., metal halides and their ion pairs, etc.) will not be discussed here, since these type of photoinitiators are beyond the scope of the present chapter. Two types of organometallic photoinitiators for free radical vinyl polymerization are considered: (1) transi-

245

Photoinitiationof Free Radical Polymerization by Organometallic Compounds

tionmetal complexeswithligand,mainly carbon monoxII. TRANSITION METAL CARBONYL ide(metal carbonyl complexes),generally, but not invarCOMPLEXES AS PHOTOINITIATORS iably, in the presence of organic halide (e.g. CC14 or CBr4), and (2) transition metal chelates. Systems of type The primary photochemical of transition metal carbon(1) contain metal in a low, often a zero oxidation state, yls [Me(CO),] involves the scission of carbon monoxide while those in type (2) contain metalin a high oxidation(CO)and the formation of coordinated unsaturated state. species:

ezcited

monoxide to regenerate the original carbonyl Me(CO),,

The first experimental evidence with the relevant theoretical background of this type of photochemical reaction was given by Koerner Von Gustorf and Grevels [l l]. The product Me(CO),- can combine with carbon

Me(C0)n-1

+ S(M)

or it may react by the addition of an n or an electron

doner (S) (e.g., monomer) according to Blyth and Hoffman [7].

m

(M)S ...Me(C0)

(7)

n-1

The species (M)S.*.Me(CO),-t arises from Eq. (7) and reacts with suitable halides (e.g.,CC4) with a generation of free radicals:

ture MFe(C0)4 in whichthe iron atom is coordinated to the vinyldoublebond. Further, ithasbeendemonstrated by the same authors that irradiation of Fe(C0h in the presence of a vinyl monomer and an organic halide leads to photoinitiated polymerization in a suitable system. The following reaction scheme is. suggested for methyl methacrylate monomer andCHBr3halide as coinitiators:

Thus, a mixture of simple carbonyls Me(CO), and halides should behave as a photoinitiator of free radical polymerization. Manysuch systems have been foundto function in this way. Complexes formed by irradiation of Fe(C0)S inthe presence of a vinyl monomer(M) (such as MMA, styrene, vinyl acetate, propylene, and vinyl ether) have been studied by Koerner Von Grustrof and colleagues [12,13] and shown to have the chemical struc-

Fe(CO)5

+ MMA

FeBr2 t 4CO .t 2 CHBr2

V

CHBr 3

H, /H C (CO),Fe+- 11

CHBr,

l

CHBT 3

-=+

(CO),Fe

r 4

L

HSC’ ‘COOCH3 Strohmeier and Hartmann [l41 first reported in 1964 the photoinitiation of polymerization of ethyl acrylate by several transition metal carbonyls in the presence of CC4. Vinyl chloride has also been polymerized in a No detailed photoinitiation mechsimilar manner [15,16] anisms were discussed, but itseems most likelythat photoinitiation proceeds by the route shown in reaction Scheme (9). The pioneering work of Bamford and coworkers in 1965[ 171and 1966 [ 181has given a detailed kineticsstudy and a mechanism of the photoinitiation of polymeriza-

\

Br

COOCH,

+

CH,=C! \

tion of methyl methacrylate (MMA) by manganese and ruthenium carbonyls in the presence of carbon tetrachloride as the coinitiator. The long wavelength limits of absorption by these materials are approximately 460 nm and 380 nm, respectively, and photoinitiationoccurs up to these wavelengths. For both photoinitiatingsystems, it was found that the rate of polymerization of MMA monomer is dependent on CC14 concentration (at constant absorbed light intensity). Quantum yieldsfor initiation at 25°C for MMA polymerization in the presence of 0.1 mol/l CC14 were determined at h = 435 nm for Mn2

Aliwl

246

been discontinued. Aftereffects of thiskind are not shown by systems containing Mn2(CO),ophotoinitiator with the same monomer. According to the kinetics and spectroscopic observations, the following reaction, Scheme (lo), was suggested for the photoinitiating system composed of Mn2(CO)lo/CCLsystem:

(C0)to and h = 365 for Re2 (C0)lo. The values reported are close to unity. However, MMA polymerization photoinitiated by Re2 (CO)lo( h = 365 nm) shows a long-lived aftereffect, persisting for several hours at 25°C after irradiation has

Mn,(CO),,

+ hv

a Mn(C0)

+ Mn(C0) 6

Reaction 10e is relatively slow theinRez(CO)lo initiating system, and the thermal reaction between Re(C0)6 formed in 10a and CC4 generated 'CC13 radicals thermally inthe dark and so is responsiblefor the aftereffect. However, Mn2(CO)lo reacts rapidly according to 10e and no aftereffect is observed for the Mn2(CO)lo/CCl, photoinitiating system. Bamford andcoworkers [l91 have shown that a prolonged aftereffect can be obtained with Mn-carbonyl in the presence of certain additives, notably cyclohexane and acetylacetone (S). It was suggested that the photochemical reaction between Mnz(CO)lo and(S) produces the active species (Z), which generates free radicals by interaction with halide andZ probably formed fromMn (CO)6 species:

Osmium carbonyl (OS,(CO)~~) acts as a photoinitiator ofvinylpolymerization[20],whichcanfunction without a halide additive.The mechanism of photoinitiation is bya hydrogen abstraction from monomerto pho-

Mn(C0)

+ CF,=CF2

Bamford and Mullik [23] have also investigated a new photoinitiating system composed of Mn~(C0)lo or Rez(C0)lo with acetylene, acetylene dicarboxylic acid, diethyl fumarate, diethyl maleate, or maleic anhydride. It was concluded that the primary radical responsible

Mn*(CO),,

e

Mn,(CO),,+ - 2CO

toexcited osmium carbonyl molecule. Addition of CCl, leads to only a moderate increase in the rate of initiation, butphotoinitiationby the Os3(CO)12/CC4 system is markedly increased by the addition of dimethyl sulfoxide (DMSO) as the electron doner. This has been explained by the formation of additional complexes with e-doner, which increases the rate of reaction with stabilization of incipientosmium cations. Strohmeier and Griibel[21]have reported that some vanadyl carbonyls ofCpV(CO)4 (Cp = cyclopentadienyl) can photochemically induce the polymerization of vinyl chloride in the presence of CCb. Bamford and Mullik1221, in 1973 demonstrated that managanese carbonyl is able to photosensize ( h = 436 nm) the polymerization of tetrafluoroethylene (TFE) in the absence of any other halide. The reaction occurs readily in bulk TFE or in carbon disulfide at -93°C; thus, Mn2(C0)lo/TFEcan photosensitize the free radical polymerization of vinyl monomers such as MMA, styrene, and acrylonitrile at 25°C. TFE as halide in this system is considerably less active than CC14 and perfluorocyclohexone iscompletelyinactive.Initiation by Mn2(CO)lo/TFE appears to be the result of electron transfer from the product of photolysis of Mn2(CO)lo (e.g., Mn(C0)4) to TFE,producing first a radical anion and resulting ultimately ina species containing M& bond:

+ -

3 [(CO),Mn -- CF2"CF,

]

for initiation comes from the addition of metal carbonyl fragments (e.g., Me(C0)4or Me(C0)6)to multiple bond in the acetylinic or olefinic derivatives. The following initiating radicals photochemicallyproduced are suggested:

Photoinitiation of Radical Polymerization Free Organometallic by Compounds

247 I

(CO),Re-C-C I Mn2(CO)ro is onlyactive in the presence of acetylene derivatives, where as Re2(CO)lo is active in the presence of both olefinic and acetylene derivatives.

CFP=CFCl -

,

CF2=CHF ,

The average quantum yield of an initiation process is 0.65, which is evidently less than the quantum yield in the Mn2(CO)lo/CC4 or Re2(CO)~ophotoinitiating systems under similar conditions. Two different mechanisms were proposed for the production of the initiating radicals, either by chlorine abstraction with formation of CF2 = 'CF primary radicals, or by the addition process, such as that occumng with the Mn2(CO)l&F4 initiating system, which would leadto initiating radicals of the type (CO)5MnCF2'CF Cl (see reaction Scheme Photolysis of several arene chromium tricarbonyls (Ar Cr (CO)3)(Ar = benzene, toluene, or xylene) and

I

I

1

and

(CO),Re-C=C I

.

1

Bamford and coworkers [24] also investigated the kinetics and mechanism of free radical polymerization of bulk MMA photoinitiatedby M n z ( C 0 ) ~ or Rez(C0)to in the presence of a series of fluoro-olefins such as:

CF2=CH2 and

CHF=CH2

photoinitiation by these derivatives have been studied by Bamford and Al-Lamee [25,26]. In the presence of active halide (e.g., CBr4 or CC&)as coinitiator, the chromium carbonyl excited state interacts with the halide molecule through an electron transfer process to give the initiating radical derived from the halide molecule. Mechanisms similarto that shown in Schemes (10) and (11) are suggested, and the quantum yield of the photoinitiation process approaches unity. In the polymerization of phenyl acetylene [27] by tungsten and molybdenumhexcarbonyls, high-polymer yields wereobtained in CC4 solvent. The following reaction scheme wasproposed, which is different from that reported by Bamford and coworkers [17-201:

~~

Copolymerization of methacrylic acidwithbutadiene and isoprene was photoinitiated by Mn2(CO)l~ without any halide catalyst [28,29]. The polymerization system is accompanied bya Diels-Alder additive. Cross propagation reaction was promoted by adding triethylaluminum chloride. It has generally been concluded that the photoinitiation of polymerization bythe transition metal carbonyls/ halide system may occur by three routes; (1) electron transfer to an organic halide withrupture of C 4 1 bond, (2) electron transfer to a strong-attractingmonomer such as C2F4, probably withscission of-bond, and (3) halogen atom transfer from monomer molecule or solvent to a photoexcited metal carbonyl species. Of these, (1) is the most frequently encountered. 111.

PHOTOINITIATION OF POLYMERIZATION BY TRANSITION METAL CHELATES

Compared with other metal coordination compounds, relatively little is known about chelate photochemistry.

The photochemistry of transition metal 1,fdiketone chelate complexes has beenknown for sometime [30,3l], and their photophysical andphotochemical properties and photocatalytic activity in different chemical reactions were reviewed in 1990 by Marciniak and Buono-Core[32]. Further discussion on the photochemistry of metal chelate will not take place here since this subject is out of the scope of this chapter. Studies in the photoinitiation of polymerization by transition metal chelates probably stem from the original observations of Bamford andFerrar [33]. These workers have shown that Mn(II1) tris-(acetylacetonate); (Mn(acac)3) andMn (111) tris-(l ,l ,l-trifluoroacetyl acetonate); (Mn(facac)3)can photosensitize the free radical polymerization of MMA and styrene (in bulk and in solution) when irradiated with light of A = 365 at 25°C and also abstract hydrogen atom from hydrocarbon solvents in the absence of monomer. The initiation of polymerization isnot dependant on the nature of the monomer and the rate of photodecomposition of M n ( a ~ a c )exceeds ~ the rate of initiation andthe initiation species is the acac radical. The mechanism shown in Scheme (14) is proposed according to the kinetics and spectral observations:

Aliwi

248

"ezcited"

inactive products

.''

M

Mn(acach in the above mechanism undergoes an intramolecular photooxidation-reduction reaction arising from the ligand to metal charge transfer process (LMCT). The quantum yield of the initiation process (c$i)is quite low 8 X indicating the great stability of the chelate ring toward photolysis. However, the quantum yield of photodecomposition (&) under similar condition is 2 X IO-', which is higher than (c$i).It is clear, therefore, that not every molecule of Mn(acac)3that is decomposedinitiates polymerization; apparently, ex-

cited moleculesof the chelate may react with a monomer to form inactive products. Photoinitiation of polymerization of MMA and styrene by Mn(fa~ac)~ was also investigated, and it was shown that the mechanism of photoinitiation is different [33] from that of Mn(acac)3 andis subject to the marked solvent effect, being less efficient in benzene than in ethyl acetate solutions. The mechanismshownin Schemes (15) and (16) illustrate the photodecomposition scheme of Mn(facac)3inmonomer-ethyl acetate and monomer-benzene solutions, respectively. (C = manganese chelate complex.)

In ethyl acetate (E):

/V

*

(CE)

M 1..

........

->

n Mn (facac),

+

facacM'

(E)

II

3 Mn(facac),

C*

II

3 Mn(facac,,

exciple2

+ +

facac' facacM'

In benzene (B)

c*

................... 3

n Mn(facac)2

............... 3

II Mn(facac),

E.

+

+

facac' facacM'

exciples

If the monomer isdenoted by (M) and photoexcited chelate by (C)*,the proposed mechanism with ethyl acetate (15,16) involves formation of an exciplex (CM*), which mayrevert to (C + M) or decompose to facacCM radicals and Mn(facac)l. This alone is inadequate, since it predicates a linear relation between the rate of initiation andthe monomer concentration, and, therefore, it is believed that monomer and ethyl acetate (E) are, to some extent, interchangeable in the reaction and that E may also form as exciplex (CE)*. The corresponding reaction Scheme, Scheme (16) proposed for benzene solution is similar,the only additional reaction being the deactivation of the exciplex (CM*) by the banzene molecule (B). Quantumyields for photoinitiationandphotodecomposition are also low in Mn(facac)3(1.5 x IO-') and

are approximately equal. This is a different behavior that observed for the Mn(acach photoinitiator. Kaeriyama and Shimura [34] have reported the photoinitiation of polymerization of MMA and styrene by 12metal acetylacetonate complex. These are Mn(a~ac)~, MoOz(acac)Z, Al(acach, Cu(bzac)z, Mg(acac)Z, Co(acat)', Co(acac)3, Cr(acac)3, Zn(acac)z, Fe(acac13, Ni(acat)', and (Ti(acac)z) - TiC16. It was found that Mn(ac ~ cand ) ~ Co(acac)3are themost efficientinitiators. The intraredox reaction with production of acac radicals is proposed as a general route for the photodecomposition of these chelates. Aliwi and coworkers have investigated many vanadium (V) chelate complexes as photoinitiators for vinyl polymerization [36-431. The mixed ligand complex of chloro-oxo-bis(2,4-pentanedione)vanadium (V). VO(acat)' Cl is used as the photoinitiator of polymerization


249

ofMMA in bulk and in benzene solution (irradiation wavelength A = 365 nm) [36]. The neutron activation analysis of the polymer reveals that initiation is effected predominantly by chlorine atoms. No retardation or inhibition were detected,

inactive p r o d u c t s

py+-

M

It has been proposed that the monomer undergo an insertion reaction into the V 4 1 bond of the photoexcited chelate molecule, resulting in decomposition of the chelate in the anon-radical route. The quantum yield of initiation isalso low, approximately 2.07 X This could be increased to 0.59 or 0.125 when 0.2 M of a strong electron doner (D) such as dimethyl sulfoxide (DMSO)or pyridine (Py) isused, respectively [36,37]. According to the UV visiblespectral and conductivity changes after the addition of DMSO or Py to the

+

-

+ PY -VO(acac) N

~ 1 '+

MMA solution of the VO(acac)zC1,it has been suggested that an ion-pair complexof the type ([VO(acac);?]+Cl-} is formed, which is photosensitive at A = 365 nm (25°C) and is able to polymerize MMA monomers. Neutron activation analysis of a polymer suggests that when Py is used as the electron doner (D), the initiation proceeds through the Cl- atom, butwhen D = DMSO, both Cl' and DMSO residues are the primary radicals produced from the photoexcited ion-pair complex. The following reaction scheme is proposed:

/ADM + - *.

] C1 I ) [ D+-()( acac)2 1c1

V

[ D-VO(acac)z C]'

but the rate of photodecomposition increases linearly with an increasing monomer concentration W], and the rate of decomposition is equal to the rate of initiation at finite [M]. The following photo-redox reaction is suggested:

hv

Iv

DMSO -VO(acac)z

V

(18)

-

W

HCl + DMSO V O ( a ~ a c ) ~ + CH~SO'CH~

The polymerization of MMA photoinitiated by alkoxo-oxo-bis(8-quinolyloxo) vanadium (V) complex [VOQZOR] has also been studied [38,39].The alkyloxo radical ('OR) formed from the photodecomposition of the chelate (A = 365) nm at 25°C) was found to be the initiating species:

hydroxy radicals derivatives (e.g., CH30 c* 'CHzOH). Again the quantum yieldof initiation is rather low being approximately equal to 2.26 x whichreflects the importance of the deactivation of the photoexcited chelate molecule by usual photophysicalprocesses. The ion-pair complex formed by the interaction of hydroxobis(8-quinolyloxo) vanadium (V) [VOQzOH] V hv Iv and n-butyl amine is also effective in photoinitiation of VOQ2 OR 6 .......... . ......_.3 VOQ2 + 'OR (19) polymerization bulk in MMA of and in solution [40]. The quantum yield of initiation and polymerization deterSpintrapping (e.s.r) andCI4-labelingtechniquesmined are equal to 0.166and 35.0,respectively. Hydrowere used to study the structure of the alkyloxy radicals xyl radical ('OH) is reported to be the initiating radical producedand to show that these radicals isomerize to and the following photoreaction issuggested: V hv 1 VOQz(n-BuNH2) I+OH-"-"I.~-BUNH~

Several vanadium (V) complexes were also studied byAliwi [41]in 1988 as possible photoinitiators for the radical polymerization of MMA. These complexes are oxo-tris(ethox0) vanadium (V), VO(OCzH5)3, oxo-tris(triphenylsiloxy) vanadium (V), VO(Si(CsH5)3)3, oxo-tris(benzoy1oxo)vanadium (V), VO(OCOCsH5)3,

N

-+VOQi:+ n B u N H z

+ 'OH

(20)

n-propylthio-oxo-bis(8-quinlyoxo) vanadium (V) VOQzSCH;?CH2CH3,andoxo-tris-indenylvanadium(V). /

Aliwl

250

In no case was retardation or inhibition detected during the polymerization process; the radical responsible for the initiation step is derived from the complex

VO(0CzHs) 3

hu

VO(OC2Hs) 2

Recently, Aliwi and Abdullah [42] have investigated the photoinitiation of styrene by oxo-tris(dimethy1 dithiocarbomato) vanadium (V) (VO(S2CN(CH3)2)3using light of A = 365 nm. Spectroscopic and kineticanalyses V

ligand produced by the scission of the V+, V-S, or V - C with reduction of V(V) to V(1V) derivatives. The following photochemical reactions for primary radical generation were suggested:

+ 'OCzHs

show that initiation occurs predominantly through scission of N,N-dimethyldithiocarbomateligand with reduction of V(V)to V(IV), andVO(S&N(CH3)2)2 is the final photolytic product.

N

VO(S2CN(Cq)2) VO(S2CN(C&)2) The quantum yieldof polymerization is 6.72 andfor photoinitiation +i = 2.85 x The polystyrene produced with this initiator shows photosensitivity when irradiated with UV light (A = 280 nm). This polymer, which cames two photosensitive end groups of - SC(S) N(CH3)2,behaves as a telechelic polymer andisituseful for production of ABA block copolymer. In another type of mixedligandoxo-vanadium Schiff base complexes: chloro-oxo-bis[N(4-bromophenyl) salicylideneiminato] vanadium (V) chelate, VOL2Cl,and methoxo-oxobis[N(4-bromophenyl)salicylideneimine vanadium(V) chelate, VOL20CH3 Aliwi and Salih [43] show that these complexes are active pho-

2

+ 'SC(S)N(CH

I

0-c

3 ) ~

toinitiators for styrene and methyl methacrylate monomers. These initiators are also active in photocuring of unsaturated polyester, and VOL2Cl could be incorporated intosecondary cellulose acetate to produce a photosensitive polymer that upon irradiation in the presence of styrene a grafted and crosslinked (network) copolymer are obtained. A number of Cu(I1) amino acid chelate Cu L2(L = glutamic acidor serine) have also been found to be effective photoinitiators for the polymerization of acrylamide 144,451. The general sequence of events is outlined in the following scheme. This was postulated to account for theresults from flash photolysisof the Cu"-glutamic acid and Cu"-serine systems.

N H d C H R w CuL \

(21)

+ H2NbHR

(27)

/

6

IV. OTHER PHOTOINITIATING SYSTEMS BASEDON METALCOMPLEXES

The participation of a monomer molecule in the initiation step of polymerization has not been required in the examples described so far. Tris(thiocyanat0) tris(pyriTO complete the informationregardingphotoinitiationdine)iron(II1)complex forms a complexwithmethyl ofradicalpolymerizationbymetalcomplexes,special methacrylate [46]. By subjecting the compound to UV photoinitiation systems will be discussed and reviewed radiation, the complex decomposes to give 'SCN as the in this section. initiating radical.

Photoinitiation of Free Radical Polymerization by Organometallic Compounds

SCN

SCN

"A

"A

+

cessfuly photoinitiated by a similar type of Co3+ complex: [COL(NH&]X2(A = 470 nm) (L = N3-, SCN-aX = NO3-, SO4')

This photoinitiating system is also used for the polymerization of other vinyl monomers such as styrene (St), acrylonitrile (AN), and vinylacetate (VA). The efficiency of photoinitiation by this system follows the order:

The initiating radicals are assumed to be 'SCN, 'ONO or 'N3 free radicals. Tris oxalate-ferrate-amine anionsalt complexes have been studied as photoinitiators (A = 436nm)of acrylamide polymer [48]. In this initiating system it is proposed that the CO*' radical anion found in the primary photolytic process reacts with iodonium salt (usually diphenyl iodonium chloride salt) by an electron transfer mechanism to give photoactive initiating phenyl radicals by the following reaction machanism:

MMA%-St>AN%VA The involvement of acrylamide monomer in the initiation step is also observed in the pentamine-aqua-cobalt(II1) complex [47]. The initiating radical is derived fromthe monomer by addition of the H 2 0 molecule with a reduction of Co3+ to Co2+. (reaction Scheme [29]) Acrylmide monomer in aqueous solution was suc-

NH3 31-7 H N

NH3

NH3

NH3

'

,C.o...,,+ HzC=CHCONH2 +S"

,..-v

H3N

1

I

HS

'SCN

Hz0

x

H3N

./

NH3 t2

CO

.,

+

'

+ N S 4 + 'CH2CHCONH2"'

4NH3

-2

hv

-3

[Fe(C204),] [Fe(C204)2]

( c ~ H J +~ CO' J ,~

.--y CO,

+

+

+y'

CD.

_....

CH2'=CH I CONH 2

Ha0

C20T4

( C , H , ) ~ I * . ~ C ~ H , I + *c,H, ~

Triphenylsulfonium tetrafluoroborate [(CaHs)3S + BF4-] is used instead of diphenyl iodoniumchloride to give phenyl radical as the initiating species. Potassium [tris(oxalato) cobaltate) (111)]with diphenyliodonium chloride also has been used as the photoinitiatorof acryl-

AgN03

'NO3

+ CH2=CHCN

+

+ -C

amide. The anion radicals CO2' and C ~ 0 are 4 ~probably the initiating species in this system [49]. Silver nitrate and/or cuppric nitrate [50] can photoinitiate the polymerization of acrylonitrile in a dimethyl formamide (DMF) medium. The photoactive species is the complex formedbetween the monomer andsalt molecules:

CH-CHCN 2-

---.-AgNO3

DJfF + CHz=CHCN +====== N 03CH2CHCN

+Ai

Aliwi

252

Okimato andcoworkers [51] have introduceda new multicomponent photoinitiatingsystem composed of the metal-ion-amine CCl, for the polymerization ofMMA (A > 300 nm). The most active system is FeS04.7H20/

r

N

r

triethyamine tetramine/CCL (1122%. Conversionwas achieved for 4 h irradiation at OOC). The following scheme was reported for the production of 'CC13 initiating primary radicals:

N

r

N

Fe

(32)

In another type of multicomponent photoinitiating system, Okimato and colleage[52] have investigatedthe Fe+3-salt/saccharide(SH) system for the photoinitiation

Fe+3+ SH

e

The following order of decreasing efficiencies for different saccharides in the svstem was found: glucose > fructose > lactose = maltose. The triplet photoexcited state of the complexfound

of acrylonitrile and acrylamide monomers in aqueous solution (A > 300 nm). The initiating radical is derived from saccharide molecule:

SH] _hvb Fg2 t

AA

+

+

H

between MMA and NiC12in N,N-dimethylaniline solvent is rewonsible for the formation of ("A) initiating radicals pioduced by H abstraction from MMA monomer (or solvent RH) by the triplet excited state molecule W].

'OH

Both 'R and MMA' radicals are found to be responsible for the photoinitiation process. Chaturvedi and coworkers [54,55] introduced phenyl dimethyl sulfoniumylide cupric chloride and chromium thiophene carboxylate as the photoinitiator of styrene and MMA. No reaction mechanism was givenfor these systems. Iwai and coworkers [56] have introduced a novel type of multicomponentphotoinitiating system for water-soluble monomer (acrylamide,acrylic acid, acrylonitrile, etc). This photoinitiator (or photocatalyst) is composed of an electron relay system in which ruthenium bipyridyl

'S

t

chloride (Ru(BPY)3C12) is acting as the photosensitizer for the generation of triethanol amine (TEA) radical cation (TEA)' by electron transfer from the amine molecule to the photoexcited ruthenium complex. The reduced formof this complex{(RU(BPY)~+} interacts with the acrylamidemonomer(AA) to give the initiating monomer radical anion (AA:) and, therefore, regenerates the original [RU(BPY)S]~+ Sensitizer, which is not consumed through the process. The following reaction scheme shows the photosensitizing process of this electron relay initiating system:


253

It was foundthat mu(1,lO phenanthroline)3I2+ and [R~(2,2'-bipyrazine)~~+] complexescan photosenstize the polymerization acrylamidein aqueous solution [57]. A similar mechanism to that shown in Scheme (34) is suggested. However, maleic acid (MA) comonomer is used as a strong photoelectron doner in the [Ru(2,2'bipyrazine)]-acrylamide/TEA relay system. The highmolecular weight, 7.7 X lo6 &mol, of the acrylamidemaleic anhydride random copolymer is obtained.

V. PHOTOINITIATEDGRAFTING AND CROSSLINKING BY TRANSITION METAL COMPLEXES The most important practical application of the organometalliccomplex photoinitiators is the possibility of using these types of initiators in modifyingthe pre-existing polymer chain, e.g., block, graft, and crosslinked copolymers preparation.

It is clear from the preceding discussion that organometallic photoinitiators (metal carbonyl or chelate derivatives) can provide a convenient route for synthesizing vinyl polymers with a variety of different reactive end group or photoreactive pendant groups or side chains through the polymer chain. In the first type, a block AB or ABA copolymer, and in the second, a grafted or crosslinked copolymer could be produced under suitable conditions. Bamford [58] first reported that a block copolymer could be producedfrom the polymerchain A carrying CC13 end groups prepared by the Mn2(CO)lo/CBr4 photoinitiating system (see reaction Schemes [lo] and [l l]) and vinyl monomer (B)in the presence of the Mnz (CO)]ophotoinitiator. Here CBr3 ends in the macro molecules and is considered a halide coinitiator. Three types of block copolymer could be formed by this technique: (1) block copolymer type ABA, when the CBr3 active groups located at one end of homopolymer (A) and the termination is exclusively by combination:

hv

-

b - A C B r 2 (B),CBr,-A

termination

'IABA" block copolymer

(2) if termination is mainly bydisproportionationor the chain transfer process is significant, AB block copolymer is produced:

2

+ACBr(B) 2

disproportimtion

n-l

B' c.................................... 3 OT

(3) ifthe initial polymercarries two terminal CBr3 groups per chain, the first product is: AB block copolymer; Br3C ACBr2 (B) CBr2 A C(Br)3. Further reactions will activate CBr3 groups in this copolymer, so that the next product will be the seven block copolymer having four blocks of A and three blocks of B and so on. Generally, therefore, the product will consist of chains of alternating blocks of A and B linked by CBr2 units and with CBr3 terminals. CBr2 groups are proven to be much less reactive towards Mn~(C0)~o than a CBr3 group. A grafted copolymer will not be considerable or not produced at all[58]. Moreover, since all primary radicals produced are macroradicals, the synthesis produces only a minimal (or nil) amount of B homopolymer (homopolymer of B is produced only by chain transfer process). Using the above method, Bamfordand Han [59] have succeeded in synthesizing and characterizing copolymers in which each block is an alternating copolymer. The block copolymers prepared are of general structure:

-

- -

poly(A-alt-B-)block-poly(C-alt-D)block-poly(A-alt-B)

(M?

2 \---ACBrB

+

2 n "AB" block copolymer

(36)

in which A = styrene,B = methylacrylate, C = isoprene, D = MMA. or A = isoprene, B = methylacrylate, C = styrene, D = MMA. or A = butadine, B = methylacrylate, C = styrene, D = MMA. Using the same method a block copolymerof polypeptides andvinylmonomerswas also prepared. As mentioned in Section 11, Bamford and Mullik [22] introduced an interesting method of photoinitiation of vinyl monomers by the Mnz(CO)loor the Re2(CO)lo/C2F4 system. By these methods polymeric molecule with (CO)sMn-CF2CF~"terminals produced is (see Scheme [12]). Ifa polymer of this kindis heated to 100°C in the presence of vinyl monomer, a block copolymer AB or ABA with CF&CF2 linkageisproduced [60]:

Aiiwi

254 0

f00

(CO),MnCF,CF,-A

\ -

c

1-3

Among the pairs of monomers used in synthesizing block copolymers by this method are: MMNAN, CzF4/ MMA, andHEMA/DMA (HEMA = hydroxyethyl methacrylate and DMA = decyl methacrylate). The block copolymer produced by Bamford’s metal carbonyYhalide-terminated polymers photoinitiating systems are, therefore, more versatile than those based on anionic polymerization, since a wide rangeof monomers may be incorporated into the block. Although the is controllablethrough the parameters mean block length that normally determine the mean kinetic chain length in a free radical polymerization, the molecular weight distributions are, of course, much broader than with ionic polymerization and the polymers are, therefore, less well defined.

Mn(C0)

,+ ‘CF, -CF,A

L ”

(37)

It is also important to note that this procedure of synthesizing block copolymers possesses an advantage in that homopolymers are avoided, since allthe initiating radicals are attached to a polymer chain. This is not so with many free radical syntheses. An analogous series of reaction gives rise to graft copolymer when the copolymer initiator carries a side chainwith active halidegroups(e.g.,CBr,, CCb, CHClCOOH or CFzCF, etc.) photoinitiation of vinyl monomer (M) by this type of a preformed halide-containing polymer in the presence of metal carbonyls (e.g., Mn2(CO),,,) then leads to synthesis of a graft anda crosslinked polymer, or a mixture of the two may be formed depending onthe character of termination of M’radicals. Reaction Scheme (38) represents the formation of a crosslink in a typical case:

CBr3

CBr3

CBr,

CBr, combination

CBr, I

A wide rangeof polymer networks are constructed in this manner. Poly(vinyltrich1oacetate) was usedas the coinitiatorwith styrene, MMA and chloroprene as crosslinking units. Polycarbonates, polystyrene, N-halogenated polyamide, polypeptides, and cellulose acetate, suitably functionalized, have been used as a coinitiator

CBr2 l

with metalcarbonyl to produce polymeric networks that are not readily synthesized by other routes. In an attempt to formulate a new photoresistant and presensitized lithographic plate, Wagner and Purbrick [61] have used poly(viny1trichloroacetate) and styrene, which together with manganesecarbonyl or phenyl chro-

Photoinitiation of Radical Polymerization Free Organometallic by

Compounds

phototoresistant or lithographic printing materials and the use of the liquid monomer (styrene) is undesirable in this technique. A polymer was, therefore, designed to overcome these shortcomings. Its chains with an average composition shown below carried both CBr3 group and polymerizable metharylate units [61].

miumtricarbonyl ( C ~ H ~ C I ~ C Owere ) ~ ) , coated onto grained anodized aluminum foil. On exposure to light, the illuminated portions were crosslinked and insolubilized so that an image developed ina mixture of ethanol and cyclohexanone was formed. Unfortunately, poly(viny1trichloroacetate)does not have suitable mechanical properties for practical use in

-

3

1.o

-

0.1

0.6

COOCH,

COOBr,

CH 3 CH2--6"-+

CH3 CH2-~6"-jf

CH2-C -----)"f

255

COOC€&

I

COOCH,

I

CH2OH

I

CH,OCOCH

CH,OOC=C% CH 3

~

~~~

Bamford and Mullik [62] have succeeded in photografting a vinyl monomer onto a styrene-MMA copolymer using the Mn2(CO),o/C2F4 photoinitiating system in acetic acid. The following scheme was reported for this process:

This system gave results that were further improved upon bythe addition of a photosensitizer (e.g., a thiopyrilium, pyrilium, and selena-pyrilium salt) to bring the absorption of light into the visible region.

-&

CH2-CH-CH2-C

CH I 3 I

Mn,(CO),o/C,F4

4 . -(- e CH,-CH-CH,-b n hv

CH 3 I

4 n

COOCH 3

COOCH 3 I

4CH,-CH-CH,-~

I

"+ n

CF, CF, Mn (CO)

(39)

COOCH 3

d-jl I

CF,

-W2

(M)x

Terminal bonding of a polymer chain to a surface is a subject of great academic and industrial interest. Eastmond et al. [63] used Mnz(C0)toand Rez(C0)lopho-

toinitiators to graft poly(methy1 methacrylate) onto a chlorinated glass surface. Active halogen groups were bonded to a glass surface by nonhydrolyzable links using the reactions set out in Scheme (40):

The chlorinated surface of glass then acted as a COinitiator, and when photochemicallytreated with metal carbonyl in the presence of MMAmonomer, a thin layer of 0.4 pm thickness of poly(methy1 methacrylate) was obtained. Grafting and modification of polymers have been found to have applications in the biomedical field. For example, poly(etherurethane), which has good elastomeric and often mechanical properties and a relatively high compatibility withblood, has been used in the man-

ufacture of prostheses, such as artificial hearts and arteries and extracorporeal circulatory systems. The highest possibledegree of hernocompatibilityis essential for certain applications, including small-bore artificial prostheses. Two general methodsof improving polymer hemocompatibility are known: (1) grafting of hydrophilic chains, and (2) the chemical attachment of an antiplatelet agent. In this respect, many hydrophilic and other monomers are grafted on solidpolyurethane surfaces such as

Aliwi

256

MMA, styrene, 2-hydroxy ethyl methyacrylate, 2-hydroxypropyl methacrylate, acrylic and methacrylic acids, acrylamide and N-vinyl pyrrolidone (NVP) [64]. By this process, polyurethane is partially converted to N-chloro or N-bromo derivatives by a short immersion

in a dilute aqueous sodium hypochlorite solution at room temperature andvinylmonomer to produce polyurethane-grafted copolymeras shown inScheme (41). This was subjected to irradiation with A = 365 nm at 25°C in the presence of Mnz(CO)lo.

hV

R'COONR"

The photografting process was also done in dimethylformamide (DMF) solution. Grafting to the solid polyurethane has the advantage of leaving the interior and the mechanical properties effectively unchanged. Furthermore, the inner surface of a tube, such as a prosthetic artery, may be halogenated by allowing the hypochlorite solution to flow through the tube for a short time; grafting is subsequently confined to the halogenated surface. Polymers withan antiplateletagent of 5-(6-carboxylhexyl) 1-(3-cyclohexyl-3-hydroxyropyl)hydantoin as an end groupwere synthesized by esterifyingthe hydantoin molecule with a haloalcohol, such as CC13CH20Hor BrCH2CH20H, and usingthe product as coinitiatorwith

VOQZOH

Mnz(C0)lo in a simple photoinitiatedfree radical polymerization. The antiplatelet activity of these polymers were compared withthat of poly(ether-Wrethane) canying the hydantoin residues in side chains. Metal chelates have also been usedin photografting and crosslinking of different types of polymers [61,65-671.

A novel type of grafting process was developed using a new photosensitive polymer containing vanadium (V) chelates. These polymers were generally synthesized by the condensation of a VOQ20H complex and a hydroxy-containing polymerto produce photoactive polymer (red in color) with pendant vanadium (V) chelate.

+ 0

OH

&Q 0 Hydroxy-containing polymerssuch as poly(methy1methacrylate-co-hydroxyethylmethacrylate) [65,66]or secondary cellulose acetate [67,68]were used for this purpose. Vanadium (V) 8-hydroxy quinoline-hydroxyethyl methacrylate adduct, prepared by condensation of the latter with a VOQ20H complex, is polymerized to

produce a photoactive polymer similarto that shown in Scheme (42). When these polymers are subjected to light of A = 365 nminbulkvinylmonomer, (MMA or styrene) grafted or extensive crosslinking polymers were produced. The photografting or photocrosslinking occurs through the macro-radicals photochemically generated on the backbone of the polymer: IV

+VoQ2 0

0

(43)

Q ~ Q 0

grafted or cross-linked

This grafting technique proved to possess several advantages: (1) the initiating macroradicals [shown in

Scheme (43)]have high activity and no other initiating species, which rise to homopolymers, are formed (i.e.,


257

no chance for any homopolymer formation); (2) unreacted vanadium chelate attached to the polymer may

be readily removedafter polymerization bysubstitution reaction with methanol or ethanol:

m+ /"\$\$\J CH30H

+

VOQ,OCH

OH

0

(3)The low quantum yield of the photografting process (@ = 2 X provides a good opportunity to control

the network formation (curing time control), and accordingly, the desirable properties of the crosslinked or grafted copolymer might be obtained. Grafting of cotton with a styrene monomeris also possible after conversion of the cotton to a secondary

cellulose acetate and after grafting, hydrolysis of acetate groups to cellulosegraftedwith styrene is possible[67]. A similar type of condensation between a hydroxylcontainingpolymer(such as secondary cellulose acetate) with VO(BrCsH4N=CH0 C6H4)2 Cl[VOL2CI] produces photoactive polymers[68].When irradiated with UV lightin the presence of styrene or "A, grafted and crosslinked polymers were obtained:

+ VO$Cl e OH

Photocuring of commercial unsaturated polyester-styrene mixture was effectively done in the presence of the VOL2CI photoinitiator complex. The chlorine atom produced by the scission of V 4 1 bond in the VOL2C1 complex is proven to be the initiating species for the photocuring process [68]. Polypyridineruthenium (11) chelate complex [Ru (BPY)~]*'is known to participate in a photoredox reaction on excitation with visible light, coupled with the +2

+

+

HCI

0 -

graftedorcross-linked copolymer

[Ru(BPy),]

(44)

SO ,,

-2

ACKNOWLEDGMENTS The author would like to thank Dr. Najat J. Salah and Mr. Saad N; Farhan for assistance in preparingand typing the manuscript.

0

tendency of the intermediateruthenium species to revert to the divalent state. Included in the reactions of these complexes are certain photoinitiating processes [56,57,69]. The reactions normally proceed via the lowest metal to ligand change transfer state, which is luminescent, and thus provides a valuablemechanistic prope. Persulfate ions are known to quench [Ru(Bp ~ ) ~ luminescence ]~' by both dynamic and static processes with the following dominant reaction:

hv

This reaction is utilized by Burrows et al. L701 to photoinitiate grafting of acrylamide, acrylic acid, methacrylamide, and acrylonitrile on cellulose triacetate in acidic aqueous solution. It was foundthat the sulfate radical anionSO4I produced photochemically in Scheme (46) is responsiblefor generating the cellulose derivative macroradicals by hydrogen abstraction, which added the vinyl monomer to produce the grafted copolymer. The main disadvantage of this method is the production of large quantities of undesirable homopolymers in addition to the grafted copolymers.

pQ ~ Q

+3

[Ru(BPy),]

+

SO,'

+

S O :

(46)

REFERENCES 1. F. C. DeSchryver, Pure Appl. Chem., 34: 213 (1973). 2. A. D. Jenkins, and A. Ledwith (eds.), Reactivity and Structure in PolymerChemistry, John Wiley & Sons, New York, Chapter 14 (1974). 3. G. Oster, and N.Yang, ChemicalReviews, 68: 125 (1968). 4. G. Oster, Encyclopedia of Polymer Scienceand Technology, John Wiley & Sons, New York, 10: 145 (1969). 5. R. S. Davidson, J . Photochem. Photobiology, A: Chem., 73: 81 (1993). 6. R. S. Davidson. J . Photochem. Photobiology, A: Chem., 69: 263 (1993). 7. J. Blyth, and A. W.Hoffmann, Ann., 53: 292 (1845). 8. D. Berthelot, and H. Gaudechon, Compt. Rend., 150: 1169 (1900). 9. I. Ostromislenski, J . Russ. Phys. Chem. Soc., 44: 204 (1912). 10. J. F. Rabek, Photochem. Photobiology, 7: 5 (1968). 11. E. Koerner Von Gustorf, and F. W. Grevels, Fortschr. Chem. Forsch., 13: 366 (1969).

258 12. E. Koerner Von Gustorf, M. C. Henry, and C. Dipietro, Z. Naturforsch, 21B: 42 (1966). 13. E. Koerner Von Gustorf, M. J. Jun, and G. 0.Schensk, 2.Naturforsch, 18B: 503 (1%3). 14. W. Strohmeier, and P. Hartmann, Z. Naturforsch, 19B: 882 (1964). 15. W. Strohmeier, and H. Grubel,2.Naturforsch, 22B: 98 (1967). 16. W. Strohmeier, and H. Grubel,Z . Naturforsch, 22B: 553 (1967). 17. C. H.Bamford, P. A. Crowe, J. Hobbs, and R.P. Wayne, Proc. Roy. Soc. A, 284: 455 (1965). 18. C. H. Bamford, P. A. Crowe, J. Hobbs, andR.P. Wayne, Proc. Roy. Soc. A , 292: 153 (1966). 19. C. H.Bamford, and J. Paprotny, Polymer, 13:208 (1972). 20. C. H. Bamford, and M. U. Mahrnud, Chem. Commun. 762 (1972). 21. W. Strohmeier, and H. Griibel, 2.Naturforsch, 22B: 98 ( 1967). 22. C. H. Bamford, and S. U. Mullik, Polymer, 14: 38 (1973). 23. C. H. Bamford,and S. U. Mullik,J . Chem. Soc., Faraday Trans I , 72: 368 (1976). 24. C. H.Bamford, S. M. Aliwi, and S. U. Mullik, J . Polym. Sci., Symp. 50: 33 (1976). 25. C.H.Bamford, and K. K. Al-Lamee, J . Chem. Soc., Faraday Trans. I , 80: 2175 (1984). 26. C.H.Bamford, and K. K. AI-Lamee, J . Chem. Soc., Faraday Trans. I , 80: 2187 (1984). 27. T. Musuda,Y . Yamamoto, and T. Higashimura, Polymer, 23: 1663 (1982). 28. C. H. Bamford, and X. Z. Han, J . Chem. Soc., Faraday Trans. I , 78: 855 (1982). 29. C. H. Barnford, and X. Z. Han, J . Chem. Soc., Faraday Trans. I , 78: 869 (1982). 30: V. Balzani, and V. Carassiti, Photo Chemistryof Coordination Compounds, Academic .Press, London Chapter 5 (1970). 31. A. W. Adamson, and P. D.Fleischauer (eds.) Concepts oflnorganic PohotochemistryJohn Wiley and Sons, London, Chapter 7 (1975). 32. B. Marciniak,and G. E.Buono-Core, J . Photochem. Photobiology, A: Chem. 52: 1 (1990). 33. C. H.Bamford, and A. N. FerrarJ . Chem. Soc. Faraday Trans. I , 68: 1243 (1972). 34. K. Kaeriyama, andY.Shimura, Makromol. Chemie,167: 129, (1973). 35. T. Okimoto, J . Polym. Sci., Part B, Polym. Letter, 12: 121 (1974). 36. S . M. Aliwi, and C. H. Bamford, J . Chem. Soc., Farady Trans. I , 70: 2092 (1974). 37. S. M. Aliwi, and C. H. Bamford, J . Chem. Soc., Farady Trans. I , 71: 52 (1975). 38. S. M. Aliwi, and C. H. Bamford, J . Chem. Soc., Farady Trans. I . 71: 1733 (1975). 39. S. M. Aliwi, and C: H. Bamford, J . Chem. Soc., Farady Trans. I , 73: 776 (1977).

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40. S. M. Aliwi and C. H. Bamford, J . Photochem. Photobi-

ology A: Chem., 47: 353 (1989). 41. S. M. Aliwi, J . Photochem. Photobiology A: Chem., 44: 179 (1988). 42. S. M. Aliwi, and S. M. Abdullah, Polym. Inter., 35: 309 (1994). 43. S. M. Aliwi, and N. J. Saleh,J . Photochem. Photobiology (in press) (1995). 44. C. N. Narnasivayan, and P. Natamjan, J . Polym. Sci., Polym. Chem. Ed., 21: 1371 (1983). 45. C. N. Namasivayan, and P. Natarajan, J . Polym. Sci., Polym. Chem. Ed., 21: 1385 (1983). 46. N. Sakota, K.Takahashi, and K.Nishihara, Makromol. Chem., 161: 17.3 (1972). 41. G. A. Delzenne, IUPAC Symp., Macromol. Preprint p. 145, Prague (1965). 48. H. Baurnann, B.Strehmel, and H. J. Tempe, Polym. Photochem., 4 : 223 (1984). 49. H.Baumann, B. Strehmel, H. J. Tempe, andU. Lammel, J . Prakt. Chem., 326: 415 (1984). 50. H.Schnecko, Chimia, 19: 113 (1%5). 51. T. Okimoto, M.Takahashi, Y . Inaki, and K. Takemoto, Angew. Makromol. Chem., 38: 81 (1974). 52. T. Okimoto, Y. Inaki, and K. Takemoto, J . Macromol. Sci., A7: 1537 (1973). 53. J. Bartbn, I. Leboc, I. Capek, and J. TKSc, Makromol. Chem., Rapid Commun. I : 7 (1980). 54. B. Chaturvedi, andA. K.Srivastava, J . Photochem. Photobiology, A : Chem., 6 4 : 183 (1992). 55. B. Chaturvedi, andA. K. Srivastava, J . Photochem. Photobiology, A : Chem., 59: 393 (1991). 56. K. Iwai, Y.Uesugi, and F. Takemura, Polym. J . 17: l005 (1985). 57. K. Iwai, M. Uesugi, T. Sakabe, C. Hazama, and F. Takemura, Polym. J., 23: 757 (1991). 58. C. H. Bamford, European Polym. J., (Suppl. 1) (1%9). 59. C. H.Barnford, and X. Z.Han, Polymer, 22: 1299 (1981). 60. C. H.Bamford, and S. U. Mullik, Polymer, 17: 94 (1976). 61. H. M. Wagner, and M. D. Purbrick, J . Photographic Sci., 29: 230 (1983). 62. C. H. Barnford, and S. U. Mullik, Polymer, 19: 948 (1978). 63. G. C. Eastmond, C. Nguyen-Hun, and W.H. Piret, Polymer, 21: 598 (1980). 64. C. H. Bamford, and I. P. Middelton, Europ. Polym. J . 19: 1027 (1983). 65. S. M. Aliwi, and C. H.Bamford, Polymer, 18: 375 (1976). 66. S. M. Aliwi, and C. H. Bamford, Polymer. 18: 382 (1976). 67. S. M. Aliwi, and B. Y. AI-Banna, Polym. Plast. Tech. Eng., 33: 503 (1994). 68. N.J.Saleh,Ph.D.Thesis,MustansiriahUniversity, Baghdad (1995). 69. C. K. Gratzel, M. Jirousek, and M. Gratzel, Langmuir, 2: 292 (1986). 70. H. Burrows, P. J. Eliseu, M. Helena, and R. M.Freire, J . Photochem. Photobipology A : Chemistry, 59: 81 (1991).

18 Chemical Modification of Polystyrenes in the Presence of Cationic Catalysis and Their Industrial Applications Ramazan Mirzaoglu, Refiga Kurbanova, and Mustafa Ersoz Selpk University, Konya, Turkey

1.

INTRODUCTION

The synthesis of new polymeric materials having complex properties has recently become of great practical importance to polymer chemistry and technology. The synthesis of new materials can be prepared by either their monomers or modification of used polymersin industry. Today, polystyrene (PS), which is widely used in industrial applications as polyolefins and polyvinylchlorides, is also used for the production of plastic materials, which are used instead of metals in technology. For this reason, it is important to synthesize different PS plastic materials. Amongthe modification ofPS, two methods can be considered, viz. physical and chemical modifications. These methods are extensively used to increase physico-mechanical properties, such as resistance to strike, air, or temperature for the synthesizing of new PS plastic materials. In the present chapter, we will start to briefly summarize all the aspects of this subject, which have been covered in previous specialized reviews, allowing for a broad and general discussion on the subject. Then we will focus on the chemical modificationof PS in the presence of cationic catalysts. When the physical modification method isused, PS is modified by mechanical stirring with various synthetic rubbers such as polybutadiene, polybutadiene styrene, polyisopropene, polychloropropene, polybutadiene styrene-acrylonitrilecopolymers. In the chemical modification, PS is modified with polyfunctional modificators in the presence of cationic catalysis. 259

New elastic polymericmaterials (resistance to higher stroke or air) can be obtained by using physical modification methods, but using this method, two phases (PS and rubber) in the mixture were formed. Small rubber particles spread as a PS layer and, after awhile, the relationship between the layers decreases and rubberparticles gather in the upper layer of the materials. This can be the cause of the loss of resistance of the materials. These material disadvantages have stimulated the polymer synthesis to increase the PS resistance to higher physico-mechanical properties, such as higher temperature and stroke for the chemical modification of PS with various functional modifiers. The reaction capability of PS is weak, but the reaction capability can be improved byanchoring the functional groupto thealiphatic chain or aromatic ring of PS using chemicalor conversion reactions. Aliphatic chain reactions are: halogenation reactions, oxidation reactions, or unsaturated acids to bonded aliphatic chain of PS (in the presence of a radical catalysis).

II. REACTION OF POLYSTYRENE A.

HalogenatlonReactions

The halogenation reaction is very important for the production of new polymeric materials [l]. If MC13 was used, the halogenation reaction took place using the aromatic ringof PS [2]. With regardto photochemical halogenation reaction is occurred either the aliphatic or aromatic ring of PS [3,4].

Mirzaoglu et al.

260

Bevington andRatti [5] reported that when CC14 was used as a solvent on the photochlorination of PS at 78"C, the reverse C atom was chlorinated. In this case, the -CC+CH,--

halogenation reaction is very important for copolymer synthesis due to theC 4 1 bond losingits resistance [6] as shown in Scheme (1).

-?-CY-

Scheme 1

B. MetallicReactions

derivatives of PS are obtained. When the derivative It is known that PS containing metal can be easily ob- compound was carboxylated, carboxyl groups (8696%) tainedbymetallic reaction. When PS solution in ben-were attached to the phenylring as L-type [7] (Scheme PI). zene is stirred withpotassium at 8WC, the potassium K(W

COOH

I

I

Scheme 2

It wasshown that chemical reactions generally occur from the aromatic ring of PS. PS has highdielectric properties andhigh hardness, but is not resistant to higher temperature and stroke and has no adhesion capability to metals. For this reason, polymeric materials having physico-mechanicalproperties could be obtained by anchoring some functional groups to the aromatic ring. For example, ion exchangeresins are prepared by amination after chloromethylation of PS and somecross-

linking polymersare also used as a membrane bydehydrochlorination. C.

Chloromethylation Reaction

Generally, PS containing amine groupsare synthesized by condensation of chlorinated PS with amines. These type of resins are widely used as anionic resins.[8] PSs containing imidazol ringshave antistatic properties and are used as additives to make dyeing of synthetic fiber materials easy [9] (Scheme [3]).

dH2CI

I

CH2-N' 'CFN

Scheme 3

ModificationofPolystyrenesinthePresenceofCationicCatalysisandTheirIndustrialApplications

D. AlkylationReaction

261

of PS with cyclohexylchloridein the presence of MC13 catalysis at 20-22°C has been reported, and the product has resistance to higher temperature of about 180-185°C ~41. The alkylation of PS with N-methylolacetamide in the presence BF3catalysis at room temperature was prepared by Swiger et al. [ 15,161who reported that the polymer has resistance to higher stroke and temperature as shown in Scheme (4).

The alkylation reaction of PS with different compounds in the presence of Lewis catalysis have been widely investigated [lo-12.Polymeric materialsthat have a resistance to high temperature can be synthesizedby bonding alkyl- or cycloalkyl groups to the aromatic ring of PS. olefins ~ O of The obtained polymer materials withC I ~ - C PS in the presence of Lewis catalysis at - 12-35°C are used as a depressing agent for oils [13].The alkylation 0 -

+

II

CH,-C-NHCbOH

L

Scheme 4

The bonding of the phosphorus group to the aromatic ringof PSis an important element of these groups, which creates a higher resistance capability to fire [17-211. The conversion of PS into its peroxide form

cannot be made possible by oxidation.This property is normally obtained from the first alkylation of PS with isopropylchloride and then oxidation of the obtained product [22],as shown on Scheme (5).

Cl C&-

CH-CCH,

C&-C-0-OH

I

CH3

Scheme 5 ~

These combination polymers can be used as a starting substance for methyl methacrylate polymerization. Polymers containing epoxy groups have been used for polymer modification, either the synthesis of branched copolymers or their monomer having amine groups [23].The PS containing epoxy group was first synthesized by Sembay et al. [24],who used styrene with epoxyperoxides (1,2-epoxy-3-tert-butylperoxypropan). However, this process is difficult and only about 0.16-0.54% of the epoxy group can be anchored.

E. AcylationReaction PS has apolar characteristics and, thus, it is difficult to or polar materials. The adhesion form a bond with metals capability of saturated polyhydrocarbonsare dependent on the basis of polar properties of polymers [25].Mitsuaki and Masyasu [26]investigated the chemical modification of PS for anchoring of the carboxyl group to PS macromolecules with maleic anhydride (MA) inthe presence of radical catalysis at 90-150°C. These authors

Mirzaoglu et al.

262

showed that bromulated PS is useful as a starting material for achieving more bonding of MA to the chain. The synthesis of acylated PS and its condensation with alde-

hydes fromthe reaction of PS withacetic acids chloroanhydride inthe presence of dcb catalysis has been synthesized in Japan patents [27,28], as shown on Scheme (6).

COCH=CH-R

COCb

Scheme 6

Therefore, the obtained -CO"CH=CH-R group PS has a high photosensitivity property and can be crosslinked with the effect of light. The group is used in the field of photography [29,30].

F. Li-Polystyrene

In American patents [31,32, the acylation of PS with acryl, methacrylate chloride, and a and P-halogen pro-. penylchloride has been done in the presence of A l a 3 catalysis at 60-100°C. This process can be accomplished by using two methods in the presence of nonsolvent; either C H d H C O C l , C I C H & H d O C l andBF3 catalysis passed from PS as vapor or acylated material

passed from a mixture of PS and AlC13 as vapor. This method is not available due to the use of more catalysis (corresponding catalysis: PS ratio is 0.4-1/1 mol). The reaction of chloromethylated PS with furanacrylic acid or unsaturated aromatic acids are preferred for the anchoring of the carboxyl group to the aromatic ring of PS. The incorporation of various functional groups to the aromatic ring of PS has been given an American patent [33].This incorporation is possible through the alkylation and acylation reaction of Li-PS with organic solvents. For instance, the various reactions are given in Scheme 7 (carbonylated PS with carboxyl acids; hydroxylated PS with ketones; epoxylated PS with epichlorohydrin).

@

CO-R

ClCO-R

" " R

I R-"-OH

CO-R

I

R'

Scheme 7

ModificationofPolystyrenes

in thePresence ofCationicCatalysisandTheirindustrialApplications

The above methods occurred in 3 steps, therefore, these methods are not preferred. For instance, in the first step, o-, m-,andp-bromostyrene and its copolymer are synthesized. In the second step, Li-PS is synthesized from the reaction of copolymers with an organic compound containingLi. The abovementioned reactions are made with different compounds of Li-PS in the third step. These methods were also investigated by Ayres and Mann [34], who used the synthesis of PS containing chloro groups withchloromethylatedPS as the first step. In the second step, formil resin wasobtained by oxidation of chlorometylated PS. In the third step, carboxylated PS was obtained by the oxidation of formol resin with acetic acid at 20°C for 48 h. There are some disad-

263

vantages in this process. The oxidation process could take place from the aliphatic chain of PS. G. Carboxylation Reaction

Another approach for anchoring the carboxyl groups to the aromatic ring ofPS is to use a two-step method [35]. In this method, first, the reaction of PS with 2-chlorobenzoylchloride is carried out. The product can then be followed as shown in Scheme (8). If chloro atom is as o-position on benzoyl chloride the reaction occurred as (a). If there is nosubstitute group o-position on benzoyl chloride, the preferred is reaction (b).

o +

COCl

Scheme 8

Compared to the chemical modificationreactions of PS, alkylation and acylation reactions are preferred to other reactions, such as halogenation, nitrolation, sulfonation, amination, and chloromethylation, etc. because the obtained polyfunctional PS has higher physico-mechanical properties.

tric properties under various conditions and also their industrial applications.Moreover, the structure of modified PS and the nature and quantity of functional groups to anchored aromatic rings have also investigated [36-521. A.

111.

CHEMICALMODIFICATIONREACTION OF POLYSTYRENE

The chemical modification of PS with epichlorohydrin (EC), maleic anhydride (MA), acetic anhydride (AA), butadiene, and isoprene in the presence of cationic catalysis such as AlC13, FeC13, BF3-O(C2H&, ZnC12,Tic&, and SnCL, have beenextensively studied under various conditions for the last 15 years. We have also studied their kinetics, physico-mechanical, thermal, and dielec-

Alkylation Reaction of Toluene as Model Compound with Epichlorohydrin

The mechanism of chemical modification reactions of PS were determined using toluene as a model compound with EC in the presence of BF3*O(C*H& catalyst and the kinetics and mechanism of the alkylation reaction were also determined under similar conditions [53-551. The alkylation reaction of toluene, with epichlorohydrin, underwent polymerization of EC in the presence of Lewis acid catalysis at a low temperature (273 K) as depicted in Scheme (9).

Mirzaoglu et al.

264 BF,:O

CHzCl

/CY

/CH2

I

'CH

+

I

H!(\'

Q

L BF,-O-CH-CHz-O

I CH CHzCl

I

CHZCl

' I

+

CH

n o ,

I

@

0

@/CH2

-

/CH2

I I

CHzCI

L BF~"O~CH"CH~fO-CH"CH~~-CCH"CH:,

I

I

CHZCI

CHzCI

I CH& I

l+"H-CHz+O-CH-CHz+nO-~=C~ 1 1

CH,CI

CHzCI

CHzCI

Polyepichlorohydrine (PEC) Scheme 9

eBF,-o-FH-cH&O-~H-CH~~O-C~H-CHZ CHZCI

CHzCI

@

+

CHzCI

.-

y C ~ C H z - p H - ~ C H z - C ~l ~ C H z - C HI - O - Be F 3

CH& I

CH& I

CH& I

n = 9-20

Epoxy toluene olygomer@TO) Scheme 10

The epoxy toluene oligomer was also obtained by alkylation of toluene with PEC as shown inScheme (10). The structure of the products were characterized by twodetector gel-chromatography(recractometer and UV). The direction of the alkylation reaction of toluene with EC was changed depending onthe reaction condition as follows: if the ratio of toluene-EC is 1:5 mol, the alkylation reaction is toward obtaining the PEC at 273 K and if the ratio is reverse at 333 K, the reaction undergone to obtain epoxy toluene oligomer is shown in Fig. 1. B. Alkylation of Polystyrene with Epichlorohydrin

It has been shown that the alkylation reaction can be used for the chemical modification ofPS to obtain epox-

ylated andpropylchlorhydrinatedPS. The PS was alkylated at first with EC in the presence of Lewis catalysis [36-401.In thismethod, the alkylation reaction is formed either by obtaining the EC homopolymer or alkylation of the aromatic ring of PS as shown in Scheme 11. The alkylation reaction of PS was dependent on the temperature. When temperature was decreased from 353 K to 293 K, the following changes occur while the PEC is increased, the quantity of EC and anchored propyl chlorohydringroup to aromatic ring wasdecreased from 4.34 mol to 2.77 mol. As shown in Scheme 11, epoxylated PS was synthesized by dehydrochlorination of propylchloride group PS with NaOH. In addition, the alkylation reaction of PS with EC can occur by alkylation of the aromatic ring or destruction of macromolecules (decreasing molecular weight) in the presence of Lewis catalysis as depicted in Table 1.

Table 1 Effect of Lewis Catalysis on Molecular Weight

and Quantity of the Obtained Products Quantity of functional group (mol%) Catalvsis Propylchlorohydrin Epoxy 3.75 BK.O(CzH5)z 17.6 3.53 AlCls ZnCh 26.7 3.24 SnC14 20.3 0.61 FeC13

Mn.10-4

21 .o

4.90 4.70 0.520.4 4.32 1.02

0.32

The effect of catalysis on the quantity of anchoring functional groups and destruction of macromolecules is obtained in the following order:

,""""* 6

BF3-O(C2H~)2 > AlC13 > SnC14 > FeC13 > ZnCl2 The kinetics and activation parameter of the alkylation reaction of PS and toluene as a model compound with EC in the presence of BFyO(C2Hs)2catalysis are given in Table 2. The initial rate and reaction rate constant was increased withincreasing temperature as

Table 2 Kinetic Parameters of PS and Toluene as a

Model Compound Wo.105(moa-s) Figure 1 The gel-chromatogram of the obtained product from the alkylation reactionof toluene with EC undervarious reaction conditions; T, K = 273 (l), 293 (21, 313 (3), 333 (4). The ratioof to1uene:epichlorohydrin;0.5 : 1 (5); 1 :1 (6), 2: 1 (7), and 5: 1 (8). (-) refractometer detector; (---)

UV-detector.

+

/"\

K*l@ (L/mol.s)

T, K

PS

Toluene

PS

Toluene

293 303 323

0.1 0.2 0.4

0.3 0.6 1.7

2.06 4.12 10.3

8.0 57.4 76.59 15.0 44.0

Lewis cat.

CH2"CH--CH,CI

m:k = (2@-38):1

Scheme 11

E., kJ/mol PS

Toluene

Mirzaoglu et al.

266

shown in Table 2. Thus, 2-oxy-3-chloropropyl group PS andoligoepichlorohydrin(OECH) were obtained.

sopropene fragments were chemically anchored to the macromolecule of PS, the physico-mechanicaland adhesion capabilityincreased. For that reason, the alkenylaC. AlkenylationReaction of Polystyrene tion reaction of PS also occurred byuse of diene hydrocarbons, butadiene, and isoprene in the presence of BF3-O(C2H5)2. The reaction can be formed by either the The chemical modification Of PS with diene hydrocarof the aromatic ring or obtaining the diene bons in the presence Of Lewis CatalYSiS are important alkenylation homopolymer(Scheme 12). for synthesizing of higher resistance, elasticity, and adhesion-capable polymers. When polybutadiene or polyi-

R=H;CH,

7-fcy-c=cH-cy-);; I

I

R

D. Acylation Reaction of Polystyrene with Maleic Anhydride andAcetic Anhydride

The acylation reaction of PS with organic anhydrides, such as maleic andacetic anhydrides, are very important for synthesizing polyfunctional (carbonyl-, carboxyl-, keto-, olefinic)PS. The incorporation of these groups to PS caused an increase of adhesion capability, physicomechanical properties, elasticity, and photosensitivity [41-461.

When MA was used as the reactive, the acylation

reaction occurred on both sides of MA both from the olefinic bond or anhydride group. However, the acylation reaction occurred from the anhydride group as shown in Scheme (13). Therefore, to fully understand the acylation reaction mechanism of PS with MA, toluene or ethylbenzene as model compounds has been investigated under similar conditions. The structure of the acylation reaction product of ethylbenzene with MA was determined by chromatographic, spectral, and chemical analysis. The results show that thereaction occurred by opening the anhydride groups and anchoring the aromatic ring.

Scheme 13

The acylation reaction of PS with MAby using model compoundsin the presence of Lewis catalysis can explained be as follows. First, it should be taken into considerationtlfat benzene is easily acylated with MA in the presence ofAlC13

[56]. This fact can serve as a model system for the study of PS-MA-Lewis acid systems. Results obtained allow one to present the general scheme for the side chain modification of PS in the presence of Lewis acid in the following way:

Modification of Polystyrenes in the Presenceof Cationic Catalysis and Their Industrial Applications

267

Scheme 14

In the case of use of BF3.0Et2 asthe catalyst, the reaction scheme can be suggested as follows: 0

+

BF,.OE$

+

~F~OOC-CH=CH-COJ

-

Scheme 15

The scheme consists of several stages: (1) complexformationof MA with Lewis acid(MX,), (2) addition of MA . MX, complex to the phenyl ringof PS, and (3) the break of the hydrogen atom from p - or o-position of the phenyl ring andits addition to the maleate fragment. In this scheme, complex-formation plays a significant role in the acylation reaction; It can also be suggested that in the acylation reaction charge transfer complex between MA and phenyl ring of PS:

. .

0

or between complexed MA and PS:

also can take place, which facilitate the processing of the acylation in the expected direction. Complex-formations between MA and aromatic carbohydrogens such as benzene and toluene were noted in earlier studies [57]. The results of analyses for carboxyl group contents, intrinsic viscosities, molecular weightsof functionalized PSs, and the ratio of virgin and acylated PS fragments (m:n) are presented in Table 3. Based on the data in Table 3, it is clear that modification of the PS by MA occurred in all cationic catalyst media, but at different levels. The highest carboxyl group concentration is found to occur with BF3-OEt2 and TiCL. Butthein case of TiCL, acylated PS has a relatively low value of viscosity, which indicates that the process of modification is accompanied by degradation of the main chain. Results of chemical analyses show that carboxyl group concentrations have increased from 4.5 to 20.8 mol% in the case of PS with M, = 2.5 x IO' and from 2.5 to 10.2 mol% for PS with M, = 5.5 X IO'. If one compares the molecular weights of modified PS, the highest was obtained in the case of BFyOEt2, which shows the ease of application of this catalyst in the modification reaction and hence activity of it. The degree of acylation of PS essentially depends on the type of catalyst used and the molecular weight of initial virginPS. As it is seen from the data in Table 3, virgin PS with higher molecular weight is acylated to a smaller degree in the equal conditions of modification.

Mirzaoglu

268

Table 3 Results of Characterization of Modified PS Prepared by Using Various Catalvsts

Initial PS with M, A.N.

2.5.10’

=

[v1

“COOH

Catalyst (mgKOH/g) (mol%)

(m :n)a

20.0 BFs.OEt2 16.6 AlCl3 20.8 Tic14 14.3 SnCL 7.7FeCl3 4.5ZnC12

Initial PS with M, = 5.0.10s

136 132 140 106 64 39

3.1

(dL/g)

5: 1 0.7510.2 6: 1 11: 0.5 8.7 1 5: 1 0.48 1 7: 1 22: 0.53 4.4 13: 1 0.543.4 22: 1 0.55 2.5

1

“COOH (mol%)

M. lo-’ 1.91 0.96 0.91 1.16 1.17 3.61 1.21

3.7

-

M,

(m:n)a 9 :1

-

-

4.05 1

30: 41:l

3.82

a Ratio of contents of virgin PS (m) and acylated PS (n) units in macromolecules. Condition: PSMA = 2.33, MNCatalyst = 1.0, 20°C,2 h.

In fact,as can beeasily traced fromTable 3, the activities of Lewis acids in the reactions studied areas follows: BF3.0Et2 2 TiCLAlCb > FeCb ZnCl2

F loo

(O/o)

c

8ol

> SnC14

Curves of precipitations (Fig. 2) of modified PS prepared inthe presence of various Lewis acids also prove that BFs-OEtz is the most convenient catalyst for the reaction studied. In Table 4, intrinsic viscosities and carboxyl contents of different fractions of modified PS obtained at different catalyst media are presented. As evidenced from these data, all PS samples used were acylated. As the viscosity of the fractions decreases, the carboxyl contents in the product increase. The highest degree of acylation for all fractions is observed when BF3.OEt2 is used as the catalyst. Existence of carboxyl groups in the structure of modified PS was also confirmed by IR studies (Fig. 3). IR spectra of acylated PS have characteristic bands at 1760, 1725,1555, and 1410 cm” corresponding to -v

t

2oh

40 6o

0.24 I “ “ “

0.40 I

0.32

aI

Figure 2 Curves of precipitationforfunctionalized PS prepared in the presence of (+) BFs-OEtz, (-0”) TiCL, (-A-) SnCL, and (-C-) FeCl3. F = the content of fractions (%); a = the ratioof solvent(benzene)/precipitant(methano1).

Table 4 Effect of Lewis Acid on the Intrinsic Viscosity and Degree of Acylation for the Various Fractions of Modified Block PS

Fractions I11

I

Catalyst

[via

BFs’OEtz AlCls Tic14 SnCL FeCb ZnClz

0.93 1.08 1.63 1.30 1.27 1.33

a

I1

-COOHb

2.12 2.38 1.37 1.85 1.64 1.26

V

[v1

“COOH

[v1

“COOH

[v]

“COOH

[v]

“COOH

0.64 0.86 1.15 0.83 1.04 0.92

2.77 5.26 2.27 3.12 2.27 2.12

0.51 0.55 0.86 0.47 0.77 0.61

3.12 9.10 3.33 7.14 4.54 2.77

0.50 0.32 0.44 0.35 0.50 0.34

8.70 12.52 8.71 10.26 7.14 4.02

0.40 0.27 0.30 0.27 0.40 0.26

17.82 16.15 14.32 15.15 7.71 5.55

[v]in toluene at 250°C (dug). “COOH is content of carboxyl group in the modified PS (mol%).

Modification of Polystyrenes in the Presence of Cationic Catalysis and Their Industrial Applications

of acyl fragments in the side chain, which are absent in the spectra of virgin PS. These bands are more intensive in the acylated PS prepared in the presence of BF3.0Et2 (spectra 5 ) and TIC4 (spectra 4).

pF3CH3COO~C&CO+

269

The mechanism of the acylation reaction of PS with acetic anhydride in the presence of Lewis catalysts can be considered as follows:

[BF3C&COO]-

+ CH3CO+

HACOCH3

Scheme 16

The anchoring of the CH3-COgroup to the aromatic ring of PS was confirmed by spectral and chemical methods and the (CH)(OH>-CH3 group was determined by hydrogenation with LiAIH4. Compared with the bonding groups (mol%) to aromatic ring of PS, the degree of acylation was observed when MA was used. These results was obtained by determination of kinetic parameters of PS with MA and AA under the same reaction conditions. As shown in Table 5 , if the initial rate (Wo) and rate constant (K) of the acylation reaction between MA and AA are compared, the MA is almost 10-14 times higher than AA in the presence of BF3-OEt2 catalyst. This fact is due to the stretching structure of MA and the effect of the catalyst.

IV. PHYSICO-MECHANICAL PROPERTIES OF POLYSTYRENE

2000

1800

Wovenumber

1500 (cm-1)

Figure 3 IR spectra of (1) virgin PS and (2-5) acylated PS synthesized by using various catalysts: (2) FeC13, (3) SnC14, (4) TiCl4, (5) BF3.OEt2.

The phy sico-mechanic and thermal properties of polyfunctional PSs that were obtained by the chemical modification with MA, AA, EC, butadiene, and isoprene in the presence of [BF3.0Et2, TiC14, AlC13, SnC14, FeC13, ZnCl2) catalysts have been investigated under various conditions. The adhesion and photosensitivity properties of new polyfunctional PSs were determined. They depended on the binding of functional groups to the aromatic ring, and especially increased with the binding of olefinic, epoxy, and carboxyl groups. In this study, we have investigated the physico-mechanical and thermal properties of synthesized PSs, which have different

Mirzaoglu

270

Table 5 Kinetic Parametersof Acylation Reaction of PS withMA and AA in the Presence of BFs.0Et2 Catalyst

K.104, L/mol.s

W0.105, moyL-s

E., kJ/mol

T, K

MA

AA

MA

AAAA

MA

298 323 343

15.8 21.6 28.3

1.1 2.1 3.1

12.7 17.1 22.7

0.9 1.6 2.3

11.3

functional groups, and the nature and the quantity of binding functional groups to the aromatic ring. In addition, the adhesioncapabilityand photosensitivity of polyfunctional PSs were also determined. The physico-mechanical, thermal, and adhesion properties of the synthesized polyfunctional PSsare dependent on the nature of functional groups in the aromatic ring. In this case, the following are properties of the chlorohydrin and epoxy groups: highest elasticity, resistance to strike, and adhesion properties with carboxyl and olefinics.Furthermore, the - C O " C H = C H X O O H group was provided new properties such as the photosensitivecapability.FunctionalizedPSsobtained are characterized by their high thermostability, adhesion, and photosensitivity. In the chemical modification of PS with MA, AA, EC, butadiene,and isoprene using cationic catalysis caused either destruction of macromolecules or the binding of functional groups to the aromatic ring. In general, physico-mechanical properties of polymers depend on the molecular weight. However, the physico-mechanical properties of PSs decreased in the presence of cationic catalysis, but increased in the case of the binding of functional groupsto thearomatic ring in spite of the destruction of PS. Therefore, new properties such as adhesion andphotosensitive capability increase

17.1

in the PS. The results of physico-mechanic, thermal, and adhesion capability properties of the synthesized polyfunctional PSs are given in Table 6. The physico-mechanic, thermal, and adhesion properties of functionalized PS are dependent on the anchored functional groups (Table 5) the following properties were observed: highest thermal properties in the XH(OH)CH2Cl, highest resistance against light in the - C O " C H = C H X O O H , highest adhesion capability

LLDPE. 7 ) of higher than LLDPE causes poor bubble stability but more downgauging in tubular blown film extrusion of the latter than the former. Han and coworkers [38,441 of commercial studied the elongational viscosity (m) HP LDPE samples. It has been foundthat ?)E increases with the increase in elongational rate and that a steadystate approaches only at low-elongational rates when studied with respect to time. VE is reported to increase with the broadening of MWD, which manifests the increase in strain hardening due to tensile stress while in film extrusion. The major processing difficulties in the tubular film blowing operation, i.e., bubble instability and breakage of the tubular blown bubble can be overcome with a better understanding of molecular parameters and their effect on rheological properties of polymer. The bubble instability results in nonuniform thickness of the film, and the bubble breakage does not allow higher take-up speed and limits downgaugingof the film. The improvement in the take-up ratio with respect to bubble rupture of HP LDPE samples is found with the narrowing of MWD and loweringthe degree of LCB since these molecular parameters help providea balance to shear thinning and meltelasticity (reduction in strain hardening) during uniaxial stretching thus improving the blowability of the polymer [38,44]. Kwack et al. [39] in their study compared melt rheology and tubular blown film extrusion of three commercial grades of LLDPE with HP LDPE samples. The VE of LLDPE samples showed a decreasing trend with the increase in elongation rate and the polymer melt exhibitedstrain softening behavior instead of strain hardening as was foundin HP LDPE samples. Similar observations have been made by others [56]. It is reported that tensile stress on the tubular film bubble both in machine and transverse directions (SIIFand S33F) increases with blow-up and takeup ratios for LLDPE but is relatively much slower compared with that of HP LDPE [39]. Figure 11 shows the change in ratio of SllF/S33F with respect to an increase in blow-upratio of HP LDPE and LLDPE samples indicating superior blowability of the latter compared with the former. The strain-softening behavior and slow rate of increase in Slip and S33F of LLDPE during tubular film extrusion facilitates extensive drawdown "downgauging" of the film [56]. On the contrary, bubble rupture occurs in the case of HP LDPE due to strain hardening and increase in melt tension during drawdown. Although an increase in the processing temperature brings down the elasticity and the tension of the melt in HP LDPE resin and, thereby, increases the possibility of downgauging. This can lead to instability of the bubble due to very low 7 ) ~ .

Mathur and Bhardwaj

284

14

-

12

-

10

-

fracture thus propagates preferentially throughthese regions 162,651. The increase in MW and control of SCB and their distribution thus help to improve properties such as impact strength, environmental stress crack resistance (ESCR), low temperature brittleness, puncture resistance (due to the increase in concentration of tie molecules), and regulationof lamellar thickness andits distribution. The increase inlength of SCB also decreases the crystallization of PES, and thus decreases the stiffness and increases the impact and tear resistance. The higher alpha-olefin comonomers such as 1octene give larger side branches, which disrupt the lamellar growth of crystallites [31]. More crystallites are thus joined by the molecules. These molecules in the amorphous phase, between crystallites, are responsible for higher levels of mechanical properties. This fits the observation that the proportion of higher alpha-olefins in the amorphous phase is higher than with l-butene. Also, that l-butene-based LLDPE needs more side chains to achieve a given density [31].

U m

m

e 4 v)

86’ 4-

21

I

2

I

I

3

I

I

1

I

I

5

Blow up ratio

Figure 11 S , I F / S ~ratio ~ F vs. blow up ratio for different HP LDPE and LLDPE samples. Source: Ref. 39.

The uniaxial VE increases with the increase in length of SCB in LLDPE, which is implied to the increase in stability of the bubble in film processing [56,58].

W.

FORMATION OF MICROSTRUCTURE DURING SOLIDIFICATION

While coolingthe melt to solidify ina processing operation, alignment of molecules controllingthe crystallinity and overall morphology depends on the molecular structural parameters of PES. The formation of the microstructure involves the folding of linear segments of polymer chains in an orderly mannerto form a crystalline lamellae, whichtends to organize into a spherulite structure. The SCB hinder the formation of spherulite. However, the volume of spherulite/axialites increases if the branched segments participate in their formation [59]. Heterogeneity due to MW and SCB leads to segregation of PE molecules on solidification [59-651. The low MW species are accumulated in the peripheral parts of the spherulite/axialites [63]. The low-MW segregated material is brittle due to a low concentration of interlamellar tie chains [65]and

V. EFFECT OF CHAIN BRANCHING ON SOLID STATE PROPERTIES

The length and degree of chain branching significantly affect the solid state properties of PES. In the case of LLDPE, the type and content of alpha-olefincomonomer determine the polymer properties. Several workers [30,38,66-721 have studied the structure properties correlation and reported the role of chain branching in the formation of the microstructure of PES. A marked difference in the properties of LLDPE and HP LDPE is due to the absence of LCBs in the former, whereas the properties of the latter are governed by the degree of LCBs. The difference inthe alignment of macromolecules during shearing and extensional flows in tubular film processing depends on length of SCBs in LLDPE [56]. This markedly affects the crystallization as well as the orientation of amorphous and crystalline structural units upon solidification [30,73,74] In their study on LLDPE resins containing l-butene, l-hexane, and 1-octene comonomers, Kalyon and Moy [30]found a significant variptionin their film thickness when measuredaround the circumference of tubular bubbles processed under identical conditions. The samples blown with a blow-up ratio of two, exhibited more significant variation in thickness than those prepared witha blow-up ratio of three. However, film processed at a higher blow-upratio has been foundto have less variation in thickness. Circumferential variation in the orientation of the film bubble is not foundto have any linear relationship with the content and length of SCBs. The resin containing a lower content of comonomer (SCBs) showed the highest value of orientation due to a higher degree of crystallinity [30].

y,

Performance Molecular Their Structure of Polyethylenes Relation to in

285

The density of LLDPE resin is affected by the concentration of SCBs, whereby an increase in the degree of chain branching leads to a decrease in crystallinity and density [30]. Thus, by varyingthe amount of comonomer, it is possible to produce LLDPE with densities in the range of0.900-0.940g/cc. Use of very high amounts of comonomers can lead to the reduction of density even below 0.900 g/cc. The density as low as 0.86 g/cc isreported for polymers havingC2 units in the main chain as 76 wt%, indicatinga very high amount of SCBs [34]. Therefore, the linear PES in this range of density havebeentermed as ultra low-density PE (ULDPE). ULDPEs are characterized by extremely low crystallinity and low meltingtemperature; Table 3 provides a comparison with HDPE. The low crystallinity and, in turn, the density are due to the interruptive effect of the alpha-olefin units duringcrystallization[34]. While comparing 1-butene ((24)-based polymer with 4methylpentene-l (Ca)-based polymer, it was found that the latter has lower crystallinity with correspondingly lower Tm as compared to the former. ULDPEs are reported to have higher impact strength and better heat sealability at low temperatures due to very low crystallinity and low meltingtemperature, respectively. Ultimate properties of LLDPE are superior to HP LDPE resins [75,76]. The tubular blown filmof LLDPE has higher tensile strength, elongation, and outstanding puncture resistance as compared to HPLDPE film. Mechanical properties data of LLDPE and HP LDPE blown film samplesprepared under identical processing conditions are shown in Table 2 [30]. The higher values of tensile strength at break andelongation at break of LLDPE samples compared with HP LDPE also indicate superior toughness of the former. This is in agreement with higher puncture and impact resistance of LLDPE film than HP LDPE. The high value of tear resistance in machine direction and lower value transverse in direction of HP LDPE as compared to LLDPE, which has a relatively higher value of tear resistance in the transverse direction, indicates a better balance of filmproperties of the latter. Lowering of comonomer content leads to higher tensile strength at yield and secant modulus and lower dart impact of the LLDPE. A comparison can be made from the test data shown in Table 2 of the polymers containing different contents of 1-hexene comonomer. 1-Butene-based LLDPE exhibits a higher solid drawability, i.e., greater elongation at break in

comparison to l-hexene- andI-octene-based LLDPE samples. This together with the observed greater drawability of the LLDPE resins as compared withHP LDPE is due to shorter branch length at equal crystallinity [30,77]. Under comparable processing conditions of tubular blownfilm extrusion, a moreuniformtensile strength in machine and transverse direction is achievable with LLDPE resins as compared with HP LDPE resins [39]. The longer SCBs provide a better overall balance of LLDPEs toughness properties than smaller ones. Most of the commercially availableLLDPE grades are based on 1-butene, l-hexene, and l-octane with l-butene-based having the major market share. Their order of toughness with respect to branch length is found to be c6 > C4 > C2. The toughness of LLDPE based on propylene comonomer is inferior to conventional HP LDPE. The rate of improvement in polymer properties begins to diminish at a point somewhere between c6 and CS on the comonomer scale. A decline in ultimate properties appears to occur with 1-decane (C,,,) comonomer [28]. The effect of comonomer type and increase of comonomer content (decrease in density) ondart impact strength of LLDPE film is shown in Fig. 12. To improve the end use performance and makethe processability easy, control ofMW and MWD as well as the use ofmore than one comonomer has been reported for LLDPE [28].UnionCarbide'shighMWLLDPE with broad MWD is a l-hexene-based resin, and its film provides superior (about 30-50% higher) tensile strength, puncture resistance, and dart impact strength than conventional l-hexene-based resin, but with lower tear resistance in the transverse direction. The broad MWD makes the resin processability easy on the conventional extruder. Use of 4-methylpentene-l comonomer with ethylene provides LLDPE resin have film properties (i.e., tensile strength, modulus, transverse direction tear strength, andimpact strength) superior to l-butenebased LLDPE resin as has been claimed by B.P. Chemicals. l-Butene has also been used as the second comonomer with4-methylpentene-l to tailor the properties of LLDPE resin [28]. The properties of 4-methylpentene-lbased LLDPE film are given in Table 4. The optical properties of tubular blown filmdepends greatly on the surface irregularities andthe size of crystallites domain in film, which, inturn, are dependent on

Table 3 Characteristics of ULDPE and HDPE [34]

"C Polymer ULD-PE ULD-PE HDPE

Tm,

Comonomer l-Butene 4-Methylpentene-l

-

75.2-84.2 40.0-40.8

129

0.887-0.895 0.865-0.867 0.956

Mathur and Bhardwaj

286

t

9/50 microns

- 300 0

-200

(2 C8 -100

c6 top0 g cl+ Cop0 C2C3

I

I

0.920

I

0.930

Density g /cm3 L

0.960

*

Figure 12 Dart impact (ASTM.Dl709) of different copolymers (at constantMW). Source: Ref. 31.

Table 4 Comparison of Film Properties of LLDPE, Based on l-Butene and 4-Methylpentene-l Comonomer [28]

Resinterpolymer propertya Tensile at break, MPa M Db T Db Tensile at yield, MPa M Db T Db Elongation, % M Db T Db Modulus, MPa M Db T Db Elmendorf tear, G/mil M Db T Db Dart drop F-50 at 660 mm

1

l-Butene copolymer

l-Butene/4-methylpentene-

4-Methylpentene1 copolymer

24.7 18.7

33.5 25.8

42.1 31.7

7.9 8.8

10.8 11.7

12.7 13.4

460 620

460 600

510 680

153.8 205.5

205.5 233.7

272.2 353.7

200 470

240 540

250 720

140

161

180

Extrusion conditions = 1.25 mil film blown at 2 and 1 blow up ratio using 63.5 mm, 20D smooth bore extruder, die gap 80 mil, melt temperature 193°C. MD = Machine direction; TD = transverse direction. Source: US. Industrial Chemicals Co.

e


287

the processing conditions [78]. The LLDPE film exhibits higher value of haze and lower value of gloss as compared with HP LDPE sample [30,39,79,80]. This is suggested to be due to higher scattering of light from the surface of LLDPE film being more rough than HP LDPE material [80]. In another study [31] it has been reported that haze of LLDPE is mainlya consequence of lightscattered by the spherulites. Thus, origin of poor clarity is not the same in LLDPE as in HP LDPE. In HP LDPE the haze is produced by surface defects coming from processing rheology [48,54,55]. Processingvariables, therefore, do not have the same effect on LLDPE as they do on HP LDPE. As described in an earlier section, the heterogenous catalysts (Ziegler-Natta catalyst) used for the manufacture of LLDPEs providenonuniform distribution of comonomers at intermolecularand intramolecular levels in the polymer chains. This structural distributioncorrespondingly reflects in the distribution on crystalline levels such as lamella thickness distribution. The wide lamella thickness distribution is the characteristic of LLDPEs that strongly influences its basic mechanical properties. Contrary to the LLDPE produced by the Ziegler-Natta catalyst, polymer prepared with a homogenous catalyst system (metallocene catalyst) has characteristic narrow structural distribution, narrower MWD, and very low density [8,37]. Due to lower density than conventional LLDPEs, these polymers are termed as very low-densityPES (VLDPEs) 181. As a consequence of narrow structural distribution, the polymer attains narrow lamella thickness distribution and morphology, resulting in superior properties (e.g., impact strength, transparency, etc.) than the polymer produced by the

Ziegler-Natta catalyst. A comparison of properties of resins and their extruded film samplesare given in Table 5. B.P.Chemicals [81] have also reported the new metallocene-based LLDPE (with l-hexene comonomer) having much superior properties than those produced by using the Ziegler-Natta Catalyst. VI.

EFFECT OF MOLECULAR WEIGHT AND MOLECULAR WEIGHT DISTRIBUTION

MW and MWD are very significantparameters in determining the end use performance of polymers. However, difficulty arises in ascertaining the structural properties relationship, especially for the crystalline polymers, due to the interdependent variables, i.e., crystallinity, orientation, crystal structure, processing conditions, etc., which are influenced by MW and MWD of the material. The presence of chain branches and their distribution in PE cause further complications in establishing this correlation. The different values of MW averages of theqolymer fall in an order, i.e., il;?z 1 > Mz MW > Mn; their approximate locations in a distribution curve are shown in Fig. 13. If all the polymer moleculesin a sample were of the same molecular size, then all the MW averages would be identical. However, this is not the case with commercial polymers. The MWD is the ratio of @ w : h and is calledthe polydispersity index.This largely vanes from one grade of polymer to the other, depending on the polymerization conditions and the type of catalyst used. Figure 14 shows different types of MWD for the polymers.

+

Table 5 Properties of VLDPEs Prepared with Homogeneous and Heterogeneous Catalyst System [g]

Comonomer l-Hexene catalyst

l-Butene

Resin properties MFR (dl0 min) 0.908 Density (g/cc) Tensile impact strength 2930 (Kg-cmlcm) Blown film properties* 3.8 Haze (%) 125Gloss (%) Dart >4000impact strength (kg cmlmm) Tear strength (kg/cm)MD/ TD Ultimate strength (kg/cm2) MD/TD * Melttemp

= 170°C; Die =

1.7 0.909 1132

1.S

3.7 130 2920

1.4

105

46/60

831120

36174

550l410

5801560

3701290

125 4-2.0t; Blow up ratio = 1.8; Film thickness

=

30 p Output = 20-25 kg/h.

Mathur and Bhardwaj

288

chromatography involvesseparation of molecules based on hydrodynamic size by sieving action.

A. MolecularWeightEffect I I

I

I l I l I

l

l

-

l

I

I L

High MW

Figure 13 Approximatelocationof MWD curve.

Low MW

-

MW averageson a

Being polydisperse, characterization of low-and high-MW fractions has been very usefulfor obtaining a better understanding of their role in polymer performance. The quantitative estimation of the macromolecules based on their MW is, therefore, carried out by using different types of fractionation methods such as addition of nonsolvent to a polymer solution, cooling a solution of polymer, solvent evaporation, zone melting, extraction, diffusion, or centrifugation. Gel permeation

Mol ecu I ar weight Figure 14 Representative differential weight distribution curve; -( ) relatively narrow distribution curve; (---) relatively broad distribution curve; (-.-)bimodal.

The effect of MWand MWD on the solid state properties have been extensively studied [11,12,82]. These studies have been made both on fractionated and whole polymer samples. Attempts have also been madeto correlate the solution viscosity, melt viscosity, MFI and other related parameters, which represent the MW and MWD of the polymers, with the solid state properties. Table 6 summarizes the results of various studies on effect of MW and MWD on the properties of PES. The different MW averages are found to provide better correlation with polymer properties viz., Mn on brittleness, impact strength and flowproperties; @W on tensile strength and hardness; and Mz on flex life and stiffness [97,109-1121. Most of the studies have been limited to evaluate the effect of a w and @n on the properties of polymers. In a recent study [l 131, the role of higher MW averages, i.e., Rz and Mz + 1 is reported in controlling the drop impact properties of moulded HDPE products. Three commercial mouldinggrades of HDPE used in this study had their densities, melt indexes, Tms, tensile strengths, and flexural moduli that were similar in magnitude. The MW,%n, andMWD values were also close to each other, but the performance in drop impact test was dissimilar. It was found [l131 that the sample havingthe highest valuesof %z and R z + 1 (i.e., long highMW tail) did not failthe drop impact test. This is in agreement with other studies [105,114] reporting that in crystallinepolymers, brittle fracture decreases because of the increase in interlamaller “tie” molecules connecting hard crystalline and soft amorphous phases with the increase in MW. The study on commercial HDPE samples could not provide a correlation of the izod impact test with the field performance test, i.e., drop impact resistance on moulded products [113]. It was found that the sample of highest density and lowest izodimpact strength passed the drop impact test, but other samples of lower density and higher izod impactstrength could not withstand shock loading bydrop impact and failedin brittle manner. This may be due to the fact that velocities and modes of loading vary widely in different impacttests. It has been reported that even the qualitative agreement between the different impact tests is poor because the test bars and mouldedproducts often have different orientation characteristics, particularlynear the surface [115]. A typical balance of processability and end use performance is the general requirement of polymeric resins. The studies on the different polymerfractions have provided a great support in tailoring the MW and MWD in order to achieve the required properties and eliminating the unwanted molecular species. The increase in low-

ses

Performance of Polyethylenes Molecular Structure Their in Relation to

289

Table 6 Effect of MW and MWD on the Solid State Properties of PE

When MW Property

[Ref.]

Spherulite

(-1

Density

(

Resistance indentatiodscratch to resistance

[84-861 No effect up [g71

When MWD down

B31

+

)a.b

to MW >l@

Hardness Tensile strengthc*d

(+)

ElongationC

(+)

Young’s modulus

No effect

Izod impact strength

(+l

Charpy impact strength‘

(+l

[12,87,89,90] [12,87,89,90] C891 [84,86,93] P71 Falling dart impact resistancee

(+l

Resistance to low temperature brittleness

(+l

1961 [12,91,96,98]

Abrasion resistance

(-)

Resistance to creep

(+)

(+l

[12,91,96,98]

[g71 [99-1011

(+l

No significant effect

[84,102]

1841,.

ESCR

(+)h

(+P

[12,84,97,102,103]

[104-1061

Softening point

(+)

Folding endurance of film

[89,102]

+

= increase; - = decrease. Measurement made on fractionated samples of HDPE[M]. Crystallinity decreases by increase in concentration of chain ends in low MW samples [85,86]. c Change in morphology and/or residual stresses due to MW and MWD may adversely effect the property (as reported for HDPE [87]). Results reported are of both fractionated and unfractionated samples. c Broadening of MWD at constant R n improves the tensile strength [92]. Increases with the increase in R n of LDPE [971. 8 Determined for HDPE[%l. h Less effective incase of HDPE [99,107]. Studied after removal of lowMW fraction [104-106]. j Narrowing of MWD reduces the amount of low MW reject material while spherulite formation thus refines the crystalline structure and lowers the void content, as a result ESCR improves [1081. a

MW fraction increases the melt flow,thus improving the processability but at the cost of toughness, stiffness, and stress crack resistance. In addition, the improvement in performance through narrowingthe MWD is restricted by the catalyst, the process hardware, and the process control limitations. Dow has developed a reactor grade HDPE of optimized breadth, peak, and shape of MWD

bycontrolling the chain length frequencies for thinwalled mouldedproducts with excellent balance of processability and performance [l 161. Wherein, the polymers of 40 and 60 MFI can be processed as 65 and 95 MFI materials, respectively, requiring lesser energy because of lower processing temperature and pressure and can fill the large mould cavity with ease. This could be

Mathur and Bhardwaj

290

achieved due to greater number of short chains. The increase in the number of long chains and thereby, highMW fraction couldprovide high toughness to the moulded product, and reduction in peak MW helped in balancing the properties (Fig. 15). B. MolecularWeightDistributionEffect

It is evident from the foregoing discussion that MW is the fundamental characteristic of polymer, controlling the performance properties. However, simple correlation of this molecularparameter can be misleading without taking the MWD into consideration. Control of MWD provides a proper balance of polymer performance characteristics. The effect of change in MWD on the properties of PES is given in Table 6 . It is wellunderstood that low MW species lubricate the polymer melt and thus improve the flow and avoid frozen-in-orientation byfast molecular relaxation inthe moulded product, whereas higher MW fractions control the mechanical properties in the solid [117,118]. Despite superior performance characteristics, film of high MW HDPE, therefore, requires more energy input in processing as compared with medium-MW HDPE. The high melt temperature built-up during processing of high-MW HDPE can also lead to polymer degradation. The highMW fraction also causes more anisotropy, which results in an imbalance of tear resistance in machine andtransverse direction of film. The narrowing ofMWD im-

C

proves the Performance characteristics of the polymer in solid state, but at the same time makes the melt less sensitive to shear rate and therefore the processing of high-MW polymerbecomesmoredifficult.Although some of the processing limitations of high MW can be overcome by broadening the MWD, it is at the cost of some mechanical properties. Broadening of the MWD strongly decreases the melt viscosity at high shear rates [l 19,1201, which inturn reduces the processing defects, such as melt fracture [121,122], and increases the melt strength [123-1261. The advantages of broad MWD polymer, both unimodal and bimodalare well known [127-1291, but most of the commercially availablegrades of polymers are of unimodel MWD. The limitations of high-MW unimodal polymers can be overcome by achieving bimodalMWD, providing a better balance of end use performance and processability. A reactor blend of both low- and highMW fractions during polymerizationprovides a polymer of bimodal MWD [ 1301. To get a product of very broad MWD with two distinct maxima for a 50:50 (by weight) blend, there must be a certain distance between the average MW of the low- and high-MWfractions. It has been reported that a homogenous polymer blend with broad bimodal MWD is impossible to achieve by mechanical mixing of two resins with an average MW ratio higher than 10 [131]. A typical properties comparison of the unimodal and bimodal high-MWHDPE resins and their films are given inTable 7. It can be well understoodthat

E 0

: 3

U (U

L \c

Increase Ion

%Increase in G imparts strength

in short- chains

d

.-c d r

t

L ,

MW

I

I

1

High molecules

I

Low MW

molecules

l--

L B r o a d e n e d MW0promotes processing/ properties balance Source

:

DOW

Figure 15 Typicalgelpermeationchromatogramsofpolyethylenes (----) new generation super flow resin. Source:

Ref. 116.

(HDPEs). -(

) conventional high-flow resin;


291

Table 7 Properties ofHighMW-HDPE [l321

Resin Property

MWD MWD

Bimodal

9 23 948

Flow index (121) MFR ( h & ) Density (kg/m3)

8

2 9

~~

Film" Thickness ( p ) 25 Dart (G) Elmendorf tear strength (Nlrnm) MD TD

12 420 190

25

570

310

4.6 20.1

6.9 44.8

13.9 32.0

12.0 21.2

Film extruded on 50 mm ALPINE, 4: 1 Blow up ratio, 80 mm die.

one can afford to increase the MW of high-MW fraction even more in bimodal MWD polymer than unimodal MWD polymer and thus further improve the end use performance without sacrificing the processibility, which is taken care of by low-MW fraction. The major advantage to the processing industry is the downgauging by use of bimodal MWD-HDPE resin; film thickness could be reduced from 20-25 pm to 12-15 p m without loss or improvement in dart impact strength (Fig. 16) [132]. The tear strength of the film has also been found to be less effected by orientation. Bohm et al. [l301 have reported the incorporation of comomomer (l-butene, l-hexene) in the high-MW fraction of bimodal PE during polymerization, obtaining a resin blend of low-MW homopolymer and high-MW copolymer. It was found that the long chain densityin-

creased in comparison to a unimodal MWD polymer containing these comonomers, but it decreases linearly with the increase incomonomer content. The stress crack resistance (failuretime) increased byapproximately two orders of magnitude as compared with unimodal MWD homo- and copolymersdue the SCB incorporated in polymer by alpha-olefin comonomer17). (Fig. Yeh et al. [l331 in their study reported a sharp increase in ESCR of linear PE (LLDPE) with the increase in

,Jim odal

Failure t i m e [h1

Figure 17 Correlation of failure time and melt flow rate MFR 190/5 of sharp notched bars under stress for unimodal homopolymer and copolymers, and bimodal copolymers. Source: Ref. 130. l

l ! !

Film thickness ( . U 1

Figure 16 Dartimpactstrength Source: Ref. 132.

I

25

ofHMW-HDPEfilms.

Unimodal Density resin (g/cc) W Homopolymer A CS-copolymer 0 C4-copolymer Cs-copolymer Bimodal resin

numbers see 0.940 +- 0.002

0.948 +- 0.002

Mathur and BhardwaJ

292

length of SCBs. It is understood that the concentration of the tie molecules with the increase in g w improve the ESCR [134-1361. The chain diffusionwithin the crystalline lamellae occurs [137]. Under stress, crack propagation insidethe semi-crystalline materialis hindered by fibrillation (crazing)at the crack tip. Tie molecules stabilize the fibrils provided they are fixed in the crystalline regions. To reduce rapid lateral diffusion of the moleculesthrough the crystallites, smallnumber of SCBs must be present along the main chain. A blend of low-MW PE and high-MW ethylene-alpha-olefin copolymer, therefore, provides excellent stress crack resistance. The bimodal MWD PE have thus been found to have superior properties such as stiffness, impact strength, and long-term stress crack resistance [l301 as compared with unimodal MWD PE. ACKNOWLEDGMENTS We wish to thank Dr. Shashikant and Dr. U. Dayal for helpful suggestions in preparing the manuscript. For the task of typing the manuscript we would like to thank Mr. A. D. Kharadi. The author, A. B. M. thanks Dr. R. C. Jain for his support and Dr. M. Ravindranathan for giving permission to publish this chapter. In addition, permissionwasgenerously granted by the copyright holders. REFERENCES 1. 2. 3. 4. 5.

6. 7. 8.

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20 The Crystallization of Polyethylene Under High Pressure Rui Huang, Xiongwei Zhang, and Guangxian Li Sichuan Union University, Sichuan, China

1.

INTRODUCTION

The actual experimentalmoduli ofthe polymer materials are usually about only 1% of their theoretical values [l], while the calculated theoretical moduli of many polymer materials are comparable to that of metal or fiber reinforced composites, for instance, the crystalline polyethylene (PE) and polyvinyl alcohol have their calculated Young’s moduli in the range of 200-300 GPa, surpassing the normal steel modulus of 200 GPa. This has been attributed to the limitations of the folded-chain structures, the disordered alignment of molecular chains, and other defects existing in crystalline polymers under normal processing conditions. Continuous efforts have been made to search for an effective approach to realize the optimum alignment of the polymer chain to achieve the maximum properties that polymers theoretically possess. In 1964, Wunderlich andcoworkers [2-41 reported that by crystallization from the melt at sufficiently highpressure, PE can form an extended-chain, crystal (ECC) rather than the normally observed folded-chain,that is several micrometers thick, much thicker than normal polymer crystal thickness (-10 nm). Fibers made in this way have a tensile strength of 2 GPa, which is even stronger than that of alloy steel wires [l]. The novel configurations and the outstanding physical performanceof ECC polymer materials demonstrate the tremendous potential remaining in the polymer itself and have provoked great interest from many polymer scientists and engineers. A great deal of research has been done since the middle 1960s on the structure-property relationship of polymers under high pressure, this research was mainly concentrated on crystalline polymers, such as PE, polychloro295

trifluoroethylene, polyamide,polyvinylidene fluoride, polyethylene terephthalate, etc. [2-381. The pressure condition used was often within the range of 0-700 MPa. These investigations ledto findings of some newcrystal structures and an understanding of the ECCgrowth mechanismand crystalline dynamics. A s a matter of fact, the study of the behavior of polymers under high pressure has become one of the significantfields of polymermorphologyandpolymer processing, and has formed a basis for developing a new approach, which parallels physical blending and chemical synthesis, toward the making of high-performance materials. The aim of this chapter is to review the development of high-pressure crystallizationof PE over the past three decades, based on the current literature. The following topics are discussed in this chapter: the ECC of polymers; the melting-crystallization of PE under high pressure; the high-pressure hexagonalphase of PE; thetypes of hexagonal phases; patterns of crystallization of PE under high pressure; theory of nucleation of ECC under high pressure; effect of molecular weight on the melting-crystallization of PE under high pressure; and the effect of other components on the melting-crystallization of PE under high pressure. II. THEEXTENDED-CHAIN CRYSTALS OF POLYMERS

Crystallinepolymers are pressure-sensitive materials due to their weak interchain potential. Therefore, they have different crystal structures under different conditions.

Huang et al.

296

It is well knownthat flexible polymerchains usually crystallize at atmospheric pressure in a folded-chain manner, givingrise to lamellar crystals with a thickness in the range of 10-15 nm, which is much thinner than the length of a fully extended chain. However, this is not the most thermodynamically favorable conformation, but itoccurs due to kinetic reasons. Crystallization can take place faster through chain folding where the chains needto continuously deposit only alonga fraction of their lengths, provided that the resulting crystal is still stable when supercooled. It is believed that the most thermodynamically stable conformation should correspond to a crystal with extended chains, a state that normally isnot realizable during primarycrystal growth because of a high-activation barrier to such an extendedchain deposition, whether enthalpic (large surface free energy [39]) or entropic (low-depositionprobability [40,41]), makingsuch crystal growth prohibitively slow. For example, Anderson [42]discovered that the lamellar thickness of a PE fraction of 12,000 molecular weight (MW) crystallized for up to 10 days at 128°C in vacua was approximately equal to its molecular length. A completely different mode of crystallization occurs at appropriately elevated pressures (hundreds of MPa). In 1964, Wunderlich andcoworkers [2-41 discovered that PE samples crystallized from a melt at about 500 MPa had a density up to 0.994 g/cm3, witha degree of crystallinity close to 100%, and a maximum melting point at atmospheric pressure of about 140°C. The morphology of crystals, revealed by electron microscopy of fracture surface replicas of the samples, was striated, banded, and had very thick lamellae. The spectrum of thickness extended to 3 pm, with an average value of 250 nm. The C axis direction, identified by electron diffraction and optical birefringence, was parallel to a characteristic striation on fracture surfaces. These structures were accordinglyreferred to as extended-chain crystals (ECC) comparedto theusual thinner folded-chain crystals (FCC) formed byatmospheric pressure crystallization. From the comparison between molecular length and crystal thickness distributions, determined by statistics of fracture surface of samples and fuming nitric acid degradation followed by gel-permeation chromatography (GPC) measurement of molecular weight distribution of the degraded samples, it has been found [5-91 that molecular length is usually longer than the crystal thickness of ECC in the chain axis direction, implying that molecular chains in ECClamellae also may be folded with their chain ends included in the lamellae, except that chain ends are not turned in for molecular lengths Oc3SnC1> OczSnClz> OcSnC13 > Pure PVC This shows a direct relationship of nonpolar n-alkyl groups withthe stabilizing effectiveness of the tin cowpound. The larger the number of nonpolar n-alkylgroups attached to tin, the more effective it isretarding in dehydrochlorination. In the case of metal carboxylates it may be the long n-alkyl chains, normallv incomorated to reduce volatil-

331

ization duringprocessing, and leachingout of stabilizer from the polymer duringthe service life of the product, that may be making the major contribution towards stabilization by acting as lubricants and reducing polar interactions within the polymer matrix. But since reactions of metal carboxylates with the PVC in the solid phase involve a series of complex reactions and the overall stability is also influenced by allthese reactions and their products, oversimplifications regardingthe mechanisms of stabilization should be made with a certain amount of caution. [1421 studied the degradation Jamieson and McNeill of poly(viny1acetate) and poly(viny1chloride) and compared it with the degradation of PVCPVAc blend. For the unmixed situation, hydrogen chloride evolutionfrom PVC started at a lower temperature and a faster rate than acetic acid from PVAc. For the blend, acetic acid production began concurrently with dehydrochlorination.But the dehydrochlorination rate maximumoccurred earlier than in the previous case indicating that both polymers were destabilized. This is a direct proof of the intermolecular nature of the destabilizing effect of acetate groups onchlorine atoms in PVC.The effects observed by Jamieson and McNeill were explained in terms of acid catalysis. Hydrochloric acid produced in the PVC phase diffused into the PVAc phase to catalyze the loss of acetic acid and vice-versa. The effect of hydrochloric acid on PVAc is shown in Eq. (37)and the effect of acetic acid on PVC in Eq. (38).

(37)

H mCH2"CH

=

CH - CH P./V

l

+ CH 3C00H

OAc

+ HCl

m CH,CH

,-CH -CH

C' O H

+ -CH

-CH 2

=

m

-CH3 CH-CH-

l

c1 + HCI + CH2COOH

Naqvi

332

The results of Jamieson and McNeill cannot be accounted for by the intramolecular mechanism proposed by Grassie and coworkers [136,137] for the thermal degradation behaviorof VCNAc copolymers (Eqs. [28] and [291). They can be accounted for much more convincingly by the alternative approach proposed by Naqvi based on polarinteractions within the PVC matrix. Just like in copolymers even in blends, the polar carbonyl group of PVAc intensifiesthe concentration of like-poles in the PVC matrix resulting in destabilization. Grassie andMcGuchan [l431 studied the thermal stability of acrylonitrile-vinyl chloride copolymers. They foundthe vinyl chloride structure in the copolymer was quite unstable and hydrochloric acid evolution occurred more readily in the copolymers than in PVC itself. It was thought that the neighboring acrylonitrile units exerted an activating influence onthe elimination. Acrylonitrile units were found to be more stable in the copolymer as compared to pure poly(acrylonitri1e). These observations can also be more convincingly explained onthe basis of ideas put forward by Naqvi. The nitrile group, which is highly polar, can be expected to exert a destabilizing influence in exactly the same way as the polar carbonyl group in VC/VAc copolymers. McNeill et al. [l441 studied the thermal stability of PVC blends with poly(acrylonitri1e) and found that poly(acrylonitrile) was stabilized by mixing with PVC, but PVC was destabilizedby the presence of poly(acry1onitrile). Substantial destabilization of PVC was observed in the temperature range where poly(acrylonitri1e) was stable in the blend. This is an interesting observation. Poly(acrylonitri1e) by its mere presence without undergoing any change itself was exerting a destabilizing influence on PVC. The explanation of mutual acidcatalysis put forward the in case of PVCPVAc blends cannot CH3 be appliedhere. The authors suggested that the physical state of the blend may be such that hydrochloric acid diffusion is hindered, giving increased opportunity for catalysis. Here also Naqvi’s hypothesis that introduction of polar groups into the PVC matrix would result in destabilization of the polymer by intensifying undesirable polarinteractions seem more plausible. Poly(metharylonitrile) was found to have a similar effect on the thermal stability of PVC. McNeill and Neil[145,146] studied the thermal stabilities of PVCand poly(methy1methacrylate)blends and observed that chain scission of poly(methy1methacrylate) occurred at a much lower temperature in the blend and the PVC was slightly stabilized. Although at first sight this would seem to contradict the effect of highly polar groups onchlorines in PVC putforward by Naqvi, in reality it does not. The carbonyl of the methacrylate group woulddestabilize the chlorines in PVC andcause HCI to evolve at a lower temperature. The catalytic efII fect of HCl on degradationof PVC would be prevented by the removal of HCl from the blend by reaction with the ester groups of poly(methy1methacrylate)and would

contribute to the stability of PVC. This could be the reason for the observed overall increase in the stability of PVC in the blend. Kovacic andcoworkers [l471 have studied the thermal degradation of blends of PVC and poly(a-methylstyrene-acrylonitrile) (PMSAN) over the whole compositionrange. These polymers are immiscible over the whole composition range as all the blends exhibit two Tgs corresponding to the TB of pure polymers. When temperature at 5% weight loss (TGA) istaken as a measure in PMSAN content, thermal stability is reduced up to 10% PMSAN content. With further increase in PMSAN content, the thermal stabilityof the blend goes through a minimumand then continuously increases with pure PMSAN being much more stable than PVC. These observations can also be explained on the basis used earlier to describe the effect of polar comonomers and homopolymers onthe thermal stability of PVC. Despite the fact that PVC/PMSAN blends are phase separated the polar nitrile groups of PMSAN present at the interface of the two phases exert a destabilizing influence on the PVC phase by intensifying the undesirable polar interactions within the PVC matrix. The reversal in thermal stability behavior of 20% PMSAN (increase in stability) is a bit difficult to explain. Some kind of phase inversion might be a factor contributing to this, but this is pure conjecture at this time. McNeill and Basan [l481 have studied the thermal degradation of blends of PVC with. poly (dimethylsiloxane) (PDMS). Structure of PDMS is given in Eq. (39). CH3

CH3

I

CH3

I -Si-0-Si-0-Si-0I I CH3

I I

CH3

(39) Every alternatingatom in the PDMS chain is the electronegative oxygen, which makes it a polar polymer. The authors observed that at low loadings of PDMS, PVC is destabilized. This is to be expected according to the alternative model for degradation and stabilization of PVC put forward by Naqvi. But for compositions with 50% or more PDMS, both polymers are stabilized. No possible explanation can be given for this reversal in thermal stability at this time. Somepolymershave both polar andnonpolar groups in the polymer chain. Thermalstability behavior of blends of some such polymers with PVC have also been studied. McNeill andcoworkers [l491 have studied the degradation of blends of PVC with poly (tetramethylene sebacate) (PTMS). The structure of PTMS is given in Eq. (40). 0

-0°C

0

II

-(CH2)g

“c - 0 - ( C H 2 ) 4 -

(40)

333

Structure, Stability, and Degradation of PVC

The polymer contains both polar andnonpolar groups. Polar carboxylate groups are separated by nonpolar ethylene units. Degradation studies of blends of PVCandPTMSshowed that incorporation of small amounts of PTMSinto PVC leads to significant stabilization in the initial stages of degradation, but in the presence of a large excess of PTMS,PVCis subject to marked destabilization. These observations can also be explained on the basis of ideas put forward by Naqvi. At low levels of PTMS, PTMS may be miscible with PVC andthe blends maybe homogeneous. The nonpolar long ethylene units in PTMSmay be reducing the undesirable polarconcentrations in the PVC matrix by acting as a buffer or diluent thereby imparting stabilization to PVC. At very high loadings of PTMS the nonpolar ethylene units may phase-separate in which case the mechanism of stabilization will become inoperative and destabilization may result due to tensions of the interface of the polar and nonpolar phases. Also, the HCI evolved during degradation may become entrapped in the polar PVC phasedue to poorer mobility inthe nonpolar PTMS phase. This will result in autocatalysis of the degradation process by HCI, which would also cause destabilization. A similar reversal in thermal stability due to the occurrence of phase-separation has beenobserved when chlorines in PVC have been progressively substituted by hydrogen atoms using tri-n-butyltin-hydride. This will be discussed in some detail in the next section. McNeill and Basan[l501have studied the blends of PVC with poly (ethylene adipate) (PEAD). Structure of PEAD is given in Eq. (41). 0

0

II

0

-C

- (CH2)4 -C

II

- 0 -(CH&

(41) Structure of PEAD is similar to that of PTMSexcept that the number of nonpolar ethylene units between polar carboxylate groups is smaller. The thermal degradation behavior of PVCPEAD blends was observed to be similar to that of PVCIPTMS blends and can be explained on a similar basis described earlier. McNeill and Basan[l511studied the thermal degradation of blends of PVC with bisphenol-A polycarbonate. The structure of bisphenol-A polycarbonate is given in Eq. (42). 0

& -*;o-c-o-

II CH3 (42)

As can be seen from the structure, the polymer contains both polar and nonpolar segments. The authors

observed that bisphenol-A polycarbonate has a stabilizing influenceon PVC over the whole composition range. From these results it appears that the blends are miscible over the whole compositionrange, and the observations can be explained on the basis similar to that used for PV/PTMS and PVCPEAD blends.

B. Effect of Nonpolar Groups Copolymerization of a small amount of a second monomer with vinyl chloride has been an important preventive stabilization technique for PVC. By including l-3% of ethylene or propylene in vinyl chloride polymerizations, copolymers of greatly improved thermalstability, compared with vinylchloride homopolymers of equivalent intrinsic viscosity, have been produced with very little effect onthe other properties characteristic of PVC [152-1541.The stabilization is believedto be due to the blocking of the progressive “unzipping” of the HCl from the PVC backbone by the second monomer. Inmy opinion this seemsrather implausible in view ofthe fact that during the degradation of PVC, the HC1 unzipping does not continue ad infinitum but tends to stop after the formation of about 7-14 conjugated double bonds as the gain in delocalizationenergyisnolongersignificant enough to keep the reaction going further. Hence, the presence of only 2-3% comonomer maynot be expected to exert a very significant effect on the ‘unzipping’ of HCl from the polymer backbone. On the other hand, ethylene andpropylene are nonpolarandwhen randomly dispersed inthe PVC, matrix, may exert a stabilizing influence by deintensifying the undesirable likepole interactions in the PVC matrix by acting as a buffer or diluent as postulated by Naqvi. Naqvi and Sen [l551 studied the thermal stability and thermal characteristics of PVC and cis-polybutadiene rubber (PBR) blends. The conclusions of their study may be summarized as follows: 1. PBR increases the inherent stability of PVC by acting as a nonpolar dispersant and reducingthe undesirable polar tensions within the PVC matrix. 2. PBR increases the stability of the degradation process itself byabsorbing the evolved HCI and preventingit from catalyzing further degradation. 3. Both the stabilizing mechanisms are limited by the degree of interaction andmiscibilitybetween the two unlike polar (PVC) and nonpolar (PBR) phases. 4. With increasingcontent of PBR (above 10%)the inherent tendency of the two unlike phases to separate out increases, resulting in reduced interaction and miscibility. The effectiveness of stabilization is accordingly reduced. KolawoleandOlugbemi [l561 have published a study of photo and thermal degradation of the two-phase

334

system of PVC and poly(isobuty1ene). They have shown that poly(isobuty1ene) hasa remarkable stabilizing influence on PVC. Polystyrene has also been found to exert a stabilizing influence on PVCthe inblends of two polymers [157,1581. These results may also be explained on the basis of the effect the nonpolar matrix has on the polar interactions within PVC. Braun et al. [l591 have published an article on the preparation of vinyl chloride-ethylene copolymers of varying composition by partial reductive dechlorination of PVCwith tri-n-butyltin-hydride and have reported the studies on the thermal stability of the copolymers. Naqvi [l601 has interpreted the observations of Braun and coworkers [l591 inthe light of the alternative approach to thermal degradation and stabilizationof PVC, based on polar interactions within the polymer matrix, proposed by him [134]. The vinyl chloride-ethylene copolymers obtained by reductive dechlorination of PVC showed a continuous increase in stability up to a substitution of about 25 mol% of chlorines by hydrogen. This stabilization effect cannot be explained merely on the basis of substitution of labile chlorines in PVC by more stable hydrogen, although this may make a nominal contribution. The levels of labile chlorines present as structural defects inPVC havebeenfound to be extremely low-much less than even 1 mol%. The reversal in the stability of the copolymers at about 35 mol% ofethylene sheds some doubt on the mechanism of stabilization based on blockingof the progressive unzipping of HCl from the polymer backbone and will be discussed subsequently. The observed increase in the stability of the copolymer may be explained onthe basis of the diluent effect the interdispersed nonpolar short ethylene sequences have onthe undesirable concentrations of likepoles in the PVC matrix described earlier. The TB of the copolymer would decrease with increasing ethylene content, making possiblethe relaxation of repulsive tensions betweenconcentrations of like-charges by molecular motion at lower temperatures. This may also be expected to contribute to the stabilization of the copolymer. The observed reversal in the thermal stability of the copolymer at a critical composition, which appears to be between 30 and 40mol%of ethylene, may be explained on the basis of the emergence of phase-separation betweenthe nonpolar ethylene and polar vinyl chlorideblocks.Although crystallization of the ethylene blocks in the copolymer is only observed when more than 70 mol% ethylene units are present, the possibility of phase-separationoccurring at lower contents of ethylene units cannot be excluded. Also, round about the critical copolymercomposition, the TBof the copolymer may be reducedto a level that would facilitate separation between the unlike phases by increased molecular mobility withinthe polymer matrix. As has been discussed earlier, occurrence of phase-separation in the copolymer would not only make the mechanism of stabilizationdue

Naqvi

to the dilution of polar-interactions in the PVC matrix inoperative but could result in its destabilization. After the critical composition at which phase-separation begins to appear in the copolymer matrix, the degree of phase-separation would progressively increase with increasing ethylene content until even crystallization of ethylene blocks becomes possibleat levels of more than 70 mol% of ethylene, i.e., a further phase-separation in the already separated polyethylene phase. The stability of the copolymer would be expected to decrease correspondingly, which is what is observed. The results of the studies of solution degradation of the copolymer samples, which showed all copolymers to be more stable than PVC, also strongly suggest that the physical state of the copolymers in the bulk may be playing a dominant role in the observed stabilizationand destabilization of the copolymers. These observations strongly favor the ideas put forward by Naqvi. The appearance of any appreciable degree of phaseseparation in the copolymer shouldbe reflected in different Tgs of the two phases. But Braun and coworkers [l591 observed only single Tgs in the reported range of 0-57.0mol%of ethylene in the copolymer. This may Constitute a criticism of the ideas put forward here. BowmerandTonelli[l611have also studied the thermal characteristics of the whole range of ethylenevinyl chloride copolymersprepared by partialreductive dechlorination ofPVCusing tri-n-butyltin-hydride. Naqvi [l621 has substantiated further his explanations for the thermal stability characteristics of ethylene-vinyl chloride copolymers reported by Braun et al. [l591 using the results of Bowmer and Tonelli [l611 as a basis. Bowmer and Tonelli [l611 have observed a single TB(545°C) for copolymers containing up to 14.7 mol% of ethylene and twoT,s for copolymers containing 15.7 mol%(24"C,60°C),29.3mol%(23"C,72"C), and 38.5 mol% (WC,57°C) of ethylene, respectively. The lower Tg obviously characterizes the ethylene phase and the higher T, the vinyl chloride phase. This is direct evidence of the presence of two phases in copolymers in the composition range 16-38 mol% of ethylene; that is, there is some degree of phase-separation between the nonpolar ethylene segments and the highly polar vinyl chloride segments. This isin agreement with the explanations put forward by Naqvi for the thermal stability behavior of the copolymers prepared in a similar manner, based on phase effects within the polymer matrix. Bowmer and Tonelli [l611 have also observed that the magnitude of the glass transition (AC,) increases with the ethylene content of the copolymer, goes through a maximum at about 30 mol%, and then continually decreases until no glass transition is observed at more than 80 mol%of ethylene. This may constitute further evidence in favor of the explanationsput forward by Naqvi for the thermal stability behavior of similar copolymers reported by Braun et al. [1591. Initially, with increasing content of nonpolar ethylene units in the co-

Structure, Stability, and Degradationof PVC

polymer, the TBand cohesive energy density decreases and internal plasticization of vinylchloride segments increases (the same factors contribute to increased thermal stability by dilution or reduction of undesirablepolar interactions in the vinyl chloride matrix), resulting in increasingA C,. A significantdegree of phase-separation between ethylene and vinyl chloride segments appears at about 30 mol% of ethylene and with further increase in ethylene content, AC, decreases. This is probably due to the fact that with an increasing degree of phaseseparation between the ethylene and vinyl chloride segments, the degree of plasticization of vinylchloride segments by ethylene segments isreduced. Also, the degree of order in the ethylene phase of the copolymerincreases with increasing ethylene content and finally results in the appearance of crystallizationof ethylene units when their content is morethan 60 mol%. Increasing order in the ethylene phase would result in a decrease in the number of its internal degrees of freedom; that is, in the possible modesof motion of the molecules, and, correspondingly, AC, would decrease. The reversal in AC, at about 30 mol% of ethylene observed by Bowmer and Tonelli [l611 and the reversal in thermal stability between 30 and 35 mol% of ethylene observed by Braun et al. [l591 tend to suggest that a significant degree of phase-separation occurs in their respective copolymer samples at about the same composition, especially in view of the explanations that have been put forward for the observed reversals. This shows that thecorrelations made between the thermal stability characteristics and thermal transitions of two differentsets of ethylene-vinyl chloride copolymers prepared in a similar manner are reasonably justified. They support the alternate explanations putforward by Naqvi [160], based on phaseeffects withinthe polymer matrix, for the thermal stability behavior of the copolymers.

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22 Solution State of Metal Complex Calixarenes and Polymeric Calixarenes Mustafa Yilmaz SelGuk University, Konya, Turkey

1.

COMPLEXES WITH ALKALI AND ALKALINE EARTH CATIONS

The calixarenes are almost totally insoluble in water and only partially soluble in most organic solvents [l]. The first success in demonstratingtheir complex abilities was achieved by Izatt and coworkers [2,31. Izatt, who has carried out extensive investigations on the complex behavior of crown ethers and various related types of compounds, perceiveda structural resemblance between the crown ethers, cyclodextrins, and calixarenes, and proceeded to test the latter for their ability to transport cations across a liquid membrane. Theyalso used an apparatus in which an aqueous source phase containing the host molecule (the carrier) and the cations were separated byan organic phase (e.g., chloroform) from an aqueous receiving phase. They discovered that although the calixarenes are ineffective cation carriers in neutral solution, they possess significant transport ability for Group I cations in a strongly basic solution. This is in sharp contrast to 18-crown-6 compounds, which are more effectivein a neutral rather than in a basic solution. Group I1 cations, including CaZ+, Ba2+, and Sr2+,are not effectivelytransported by these calixarenes. Control experiments withp-tert-butyl phenol, whichshows little or no transport ability itself, support the idea that the macrocyclic ring playsa critical role, although it is not yet clear what the role is exactly. The diameters oofthe annuli of the calixarene mono anions are ca. 1.0 A for the cyclic tetramer, 2.4 A for the cyclic hexamer, and 4.8 W for the cyclic octamer. Thus, for complexes in which the cation and oxygen are coplanar, the cyclic tetramer has too small an opening even for Li+ ,whereas the cyclic octamer has too large an opening to fit snugly, 339

around even CS+.The cyclic hexamer seems to be the system best constituted to behave in crown ether-like fashion and, indeed, even the synthesis of p-tert butyl calix[6]arene may be influenced by a template effect. Extending their initial studies to include multicarbon systems comprised of equimolar mixturesof two, three, and fourcations from NaOH, KOH, RbOH, and CsOH, Izatt and coworkers [3] found that selective transport of CS+ occurs in all cases. In another work in this field, Izatt and coworkers [4] have a patent that has been issued for a process that uses this technology, and was utilized to recover cesium from radioactive wastes. Bocchi et al. [5] pointed out that while the tetraacetate of p-tert-butyl calix[4]arenefails to form complexes with guanidium ionsor CS+,the octa-(3,6-dioxaheptyI)ether of p-tert-butyl calix[S]arene forms strong complexeswith these cations. The Parma group [61 later showed that the hexa-(3-oxabutyl)ether of p-tert-butyl calix[6]arene also forms complexes with these ions. An extensive survey hasbeen carried out by McKervey and coworkers [7], who prepared the carboalkoxymethyl ethers of p-tert-butyl calix[4larene, ptert-butyl calix[6]arene, p-tert-butyl calix[8]arene, calix[4]arene, calix[6]arene, and calix[8]arene, and measured their abilities to extract cations from the aqueous phase into the nonaqueous phase. They concluded the following general aspects for the phase-transfer experiments: (1) the calix[4]arene compounds show the greatest selectivity for Na+ ;(2) phase-transfer of Li+ is inefficient with all of the compounds; (3) the calix[6]arene compounds showless affhity for Na+ than for K + , with plateau selectivity for Rb+ and CS+;(4) the calix[S]arene compounds are the least efficient of the cyclic oligomers, showing low levelsof transport and low discrimination for all five cations; (5) the calix[6]arene

Yilmaz

340

compounds are significantlymoreeffective than 18crown-6 for Na+ and K + and much more so for CS+; and (6)the tert-butyl group appears to increase to some extent the selectivity of the calix[4]arene compoundsfor Na+. Calestani et al. [8]prepared dietylamide of the tetracarboxymethyl ether of p-tert-butyl calix[4]arene and reported alkali metalpicrate extraction constants of 1.9

x 1O9forNa’,2.8 X 107forK+,and1.3X 107forLi+. These constants were of a magnitude that persuaded the authors to liken the binding properties of these calixarenes to cryptands and spherands. McKervey [71 and Chang and Cho [91 carried out similar studies. They prepared various ester derivatives of calixarene and tested their ion carrying capacities (Scheme l).

a R’=tert-butyI,R=Et,n=4 b R=tert-butyl,R=Me,n=4 c R=H, R=Efn=4 d R=H, R=Me,n=4 e R’=tert-butyl, R=Et,n=6 f R=tert-butyl, R=Me,n=6

g R=H, R=Et,n=6 h R=H, R=Me n=6

i R’=tert-butyl,R=Et n=8 j R=terkbutylR=Me, , n=8

k R‘=H, R=Et, n=8 I R=H, R=Me,n=8

Scheme 1

Chang and coworkers [l01 have synthesized amide derivatives of calixarenes and examined their ion binding properties with Group I and Group I1 cations. They observed that although the amides are much less effective than the esters for the complexation of Group I cations they are more effective for Group I1 cations. Thus, it has been shownthat calix[4]aryl esters-exhibit remarkably high selectivity toward Na+ [ll-141. This is attributable to the inner size of the ionophoric cavity composed of four OCHzC=O groups, which is comparable to the ion size of Na+, and to the cone conformation that is firmly constructed on the rigid calix[4]arene platform (Scheme 2).

NaH should serve as an efficient “template ion” to yield cone conformers. Conformer distribution for the reaction ofp-tert-butyl calixr41arene andethyl bromoacetate is given in Table 1. Kimura andcoworkers [171, Diamond [181, and Damien et al. [l91 havedescribed that the polymeric calix[4larenes have been usedas ionophores in ion selective electrodes for Na+ (based on calixarene esters and amides) and for Na+ and CS+(based on p-alkylcalixarene acetates). The electrodes are stated to function as potentiometric sensors as well, having good selectivity for primary ion, virtually noresponse to divalent cations, and being stable over a wide pH range. Other recent works in this field, studies on the transport of alkali and alkalineearth cations with p-tert-butyl calix[n]arene esters and amides, were carried out by Arnaud-Neu et al. 1201 and Casnati et al. 1211. They prepared 1,3-alternate calix[4]arene-crown-6 as a new class of cesium selective ionophore. In another work, ionophoric bis-calix[4]aryl esters in which the lower edgesconfront each other were synTable 1 Conformer Distribution for the Reaction of p-tert-Butyl

Calix[4]arene and Ethyl Bromoacetate

D61 Distribution of

R== , H , Scheme 2

a

in Scheme 1

Solvent Partial-cone Cone Base 1,3-Alternate

100 THF Shinkai [l51 concludedthat p-tert-butyl calix[n]arLizC03 DMF ene tetra esters form stable monolayers at the air-water NazCO3 DMF interface and the metal responds, therein, quite differ16 DMF ently fromthat in solution. They reported 84 that examinaDMF tion of the ,metal “template effect” on the conformerNazCO3 Aceton distribution establishedthat when the metal cation pres96 KzCOs Aceton ent in the base used serves as a “template,” the coneCszC03Aceton conformer results are predominant [16]. Hence, Na+ in

NaH 100 88

CSZCOS 73

0 0

12

27

100

0

0

3 100

0 0 0 0 0 0 0 0

Solution

of Caiixarenes

341

thesized by Ohseto, et al. [22]. The calixarenes are interesting compounds for which the unique intramolecular metal-hopping is observed using the 'H-NMR spectral method. The majority of the literature reports deal with the reaction of calixareneswith Group I and I1 cations. Polymeric calixarenes have beenthe subject of a more recent innovation. Hams etal. [23] have prepared a calix[4]arene methacrylate, its polymerization,and Na+ complexation (Scheme 3). They concluded that both monomers and polymers form stable complexes with sodium thiocyanate. Recently, Deligoz and Yilmaz [24,25]described the preparation of two polymeric calix[4]arene tetra esters (Scheme 4) and their Na+-complexation. Based on phase-transfer experiments with these compounds using alkali picrates in water-dichloromethane, they confirmed that polymers are asNa+ selective as monomers. R=C&C02Et., X=CH2CO2CH2CH2OCO Scheme 3

2

3

(n-38-42)

/

R- CH,COqH, 5

Scheme 4

0

\

342

11.

Yilmaz

SELECTIVE IONOPHORES FOR METAL CATIONS

from sea water per 0.1 g of resin indays 7 at a flow rate of 30 mL/min. Another polymer-bound analog has also been reported by Shinkai et al. [36] whotreatedp-(chloThe most important contributions this in field were made rosulfonyl)calix[6]arene with polyethleneimine and obShinkai and coworkers[26-321. They have embarkedon tained a gel-like product, which contained one calixaran ambitious program that focuses on the uses to which ene unit for each 15 ethyleneimine units. It exhibited a calixarenes can be put. A good example, and one that similar binding power and selectivityfor U022 as that represents an especially interesting study of cation comof the parent calixarene, p-~ulfonatocalix[6]arene. plexation by calixarenes, deals with the extraction of Harrowfield et al. [37-391 have describedthe strucuranium from sea water. The world’s oceans contain a tures of several dimethyl sulfoxide adducts of homo bitotal of about 3 billion tons of uranium in the form of metallic complexes of rare earth metal cations with pU022+associated with carbonate. Although this repretert-butyl calix[8larene and bis-ferrocene derivatives of sents an enormous quantity of material, its concentrabridged calix[4larenes. Ludwinget al. [40] describedthe tion is only about 3 parts per billion, an amount that solvent extraction behavior of three calixarene-type corresponds to less than 1 mg in a large-size backyard cyclophanes toward trivalent lanthanides La3+ (Ln = swimming pool. Also,the U0z2+ is accompanied by nuLa, Nd, Eu, Er, andYb).Byusingp-tert-butylcamerous other cations, most of them present in a much lix[6]arenehexacarboxylic acid, the lanthanides were larger concentration. Thus, the extraction of uranium extracted from the aqueous phase at pH 2-3.5. The exfrom sea water poses a tantalizing challenge, whichhas tractability is: Nb, Eu > La > Er > Yb. been addressed by a number of chemists duringthe past Yosida et al. [411 found that p-tert-butylcalix[6]ardecade. If the “greenhouse effect”proves to be responene can extract Cu2+from the alkaline-ammonia solusible for adverse global changes in the climate, it may tion to the organic solvent. Nagasaki and Shinkai [42] force greater attentionto nuclear energy,and the recovdescribed the synthesis of carboxyl, derivativesof calixery of uranium fromsea water will become a more press[nlarenes (n = 4 and 6) and their selective extraction ing problem in the future than it is at the present time. capacity of transition metal cationsfrom aqueous phase Early reports of U02’+ complexation came from the to the organic phase.Gutsche and Nam [43] have synthelaboratories of Alberts and Cram [331in1976, where sized various substituted calix[nlarenes and examined crown ethers chemistry was coming to fruition. Another the complexes of the p-bromo benzene sulfonate of pwork in this field came from Japan under the leadership (2-aminoethyl)calix[4]arenewith Ni2+,Cu2+,Co2+,and of the late Iwao Tabushi, who found the thatmacrocyclic Fe2 . triamine-l is particularly effective [341, having a K,,,,, The titanium complexes of calixarene were obtained of for U022+. This material was tested in the by Olmstead et al. [441 and Bott et al. 1453, who examocean by placingit 10 m belowthe surface in the strong ined their x-ray characteristics. Recent research in that Kuorshio Current off the coast of Mikura, an island in field has been conducted by Rudkevichet al. [46]. They Japan, where it absorbed 50 pg of uranium per gram of preparedcalix[4]arene-triacids as receptors for lanresin per day [35]. tanides. Yilmaz and Deligoz [47,481 prepared a calix[6]arene S \c/SNa substituted with six aminoglyoxime groups (21, and two vic-dioxime compounds of calix[4larene(3,4), and they examined their chelating ability with Co2+, Cu2+,and Ni2+. Yilmaz and Deligoz[491 have theorizedthat calixarenes might be utilizedas selective ionophoresfor Fe3 . Thus, they studied the selective liquid-liquid extraction of Fe3+ cation from the aqueous phase to the organic phase by usingp-tert-butyl calix[4larene (l),calixt41arene (2), p-nitro-calix[4larene (3), calix[4larene-p-sulfonic acid (4), p(diethyl-amino)methylcalix[4larene (5). tetra-methyl-p-tert-butyl calix[4]arene-tetraketone (6), 25,27-dimethyl-26,28-di-hydroxy-p-tert-butyl calix[4lar1 ene-diketone (7), calixt41arene bearing dioxime groups on the lower rim (g), and a mono-oxime (9). The affect A polymer-bound analog of the p-sulfonatocalix[6]arenes is described in a Shinkai patent [31,32], of varying pH upon extraction ability of calixarenes substituted with electron donating and electron with drawwhich states that the hexakis(carbetoxymethy1)ether of p-sulfonatocalix[6]arene waspartially nitrated, amiing groups at their p-position. Observed results were nated, and fixedon crosslinked chloromethylated polycompared withthose found for unsubstituted calix[4larene. Compounds 3 and 4 were used as electron withstyrene. This resinis stated to absorb 108 pg of uranium +

+

+

343

Solution State of Calixarenes

2 H

\=WoH

0 0

0 0 4

3

1 : R'= c(CY),

R= H

2R: = H R=H 3:

R'=NO,

R= H

4:

R'=SO,H

R=H

5 : R' = CH,N(C,H,), 6 : R' = C(CH,),

R= H R = CH&OCY

7: R = CH,COCH,

N-OH II 8: R = CH,CCH,

OH 9

YCO

C

H

=d

Yilmaz

344

drawing and 1 and 5 as electron donating calixarenes (Table 2). The results of the extraction experiments indicated that the nature of p-substituted group has not significantly affected the extraction process. The extraction of Fe3+ ions was increased by increasing pH, but exhibited a decrease at pH 2.2; 51.0% extraction was accomplished at pH 5.4 with 6 in which all phenolic groups were substituted. When compound 7, in which two phenolic groups were substituted, was used, the extraction ratioincreasedsignificantly(Table 2). The above observations evidently indicatethe important role of phenolic oxygen in this procedure. Extraction experiments with the compound 8 yielded similar results as obtained with other compounds. The effect ofpH on the extraction of 8 was observed to be lower; 62.8% extraction was accomplishedeven at pH 2.2. This result is due to the presence of two adjacent oxime groups (-"-OH) in the compound 8. Since the extraction experiments performed with its monomer 9 the ratio was only 3.6% at pH 2.2. The above observations indicate that the cone conformation of calixarene bearing the oxime groupsalso plays an important role in the extraction process. The ultraviolet spectrum of compound 2 in DMF did not exhibit an absorption maximum above 300 nm. Formation of such a complex was evident from the changing color of solution into brown and the appearance of an absorption maximum at 534 nm. The metal-ligand ratio determined at this wavelength by the Job method was 1:1. The extraction reaction of the present systems can be expressed by Eq. (1).

Where aq and org denote the species in the aqueous and the organic phase, respectively. The extraction equilibrium constant (Kex)is given by :

log D

=

n pH

+ log Kex + log [LHmlorg

(4)

Equation (4) indicates that the slope n of the log D versus the pH plotcorresponds to the number of protons released upon extraction. If the logarithm of the ratio between Fe3+ content in aqueous phase and organic phase is plotted as a function of pH, a linear relation was obtained betweenpH 3.5-5.4, which deviated from linearity at lower pH values (2.2). The fact that the slope of the curves were very close to unity indicates that only one proton has separated from the ligand (Eq. [5]). Fe3+ + HL =SLFe2+ + H +

(5)

Since one proton separation will hardlyoccur at pH

2.2, extraction ratios of Fe3+ with compounds 1-7 are quite low. Yet compound 8, which carries oxime groups,

was able to extract a considerable amount of Fe3+ at pH 2.2. The logarithmic extraction constants logKex(Kexin moVL) corresponding to Eq. (4) are as follows: log Kex = 1.79 2 0.15 (1) log Kex = 1.50 2 0.15 (4) log K,, = 3.40 2 0.10 (5) log Kex = 1.80 2 0.10 (7) The solvent extraction mechanism with1-5 and 7,8 is different from 6, since 6 is a remarkably good ionextracting compound for Na+ ions [12-141. In a twophase solvent extraction of Fe3+ from the organic phase into the aqueous phase with 6, can be explained interms of the exchange of Fe3+ ions by the Na+ ions, which isin an organic phase. The extraction processes are shown as in Fig. (1). In order to ascertain that the Na+ salt of 6 indeed exists in the chloroform phase, an aqueous solution (25 mL)containing 0.01 M CH3COOH/ CH3COONa buffer (pH4.5) and a chloroform solution (5 mL)containing 5.3 x M of 6 were mixedand shaken. After 12 h at 25°C the chloroform phase was separated and extracted with 0.1 M HCI solution. The flame photometric analysis of this solution established that 70% of 6 is dissociated as the Na+ salt. Based onthe above results they have concluded that the ligand groups circularly arranged on the lower rim of the calixarene cavity construct a novel cyclic metal receptor for selective extraction of transition metal cations. Results suggest that the fine tuning in molecular

Table 2 Extraction of Fe3+ Cation with Ligand (%) [49]

1 22.448.3 46.5

L2

PH 2.2 3.818.5 4.5 5.4

7.0 40.5 66.0

8.4 56.0 28.2 57.5 90.0

L3

L4

L5

L6

L7

L8

L9

27.1 42.1 77.0 77.1

15.7

22.0

12.0

20.7

62.8

3.6

65.O 77.8

72.0 84.5

20.4 51.o

28.1 87.8 40.7 92.1

88.6 89.4

Aqueousphase[Metalnitrate = 1.06 X MI. organicphase, [chloroform, (ligand) = 5.3 x M].pH: 2.2 (0.01 M NaN03/HN03, with KCO, pH 3.8, 4.5 and 5.4 (0.01 M CH3COONa/CH3COOH, p = 0.1 with KCI), at 25°C for 12 h.

p = 0.1

Solution State of Calixarenes

345

Figure 1 Extraction mechanism proposed for La.

design can bedone by using functional groups arranged on the lower rim (closed side of the calixarene cavity) rather than by using those arranged on the upper rim (open side of the calixarene cavity). Recently, Deligoz and Yilmaz [51] prepared three polymeric calix[4]arenes, which were synthesized by reactingchloromethylated polystyrene with 25,26,27tribenzoyloxy-28-hydroxy calix[4]arene (Za, 3a) and POlyacryloylchloridewith 25,26,27,28-tetraacetoxy calix[4]arene (4a). After alkaline hydrolysis of the polymers,theywereutilized for selective extraction of transition metal cations from aqueous phase to organic phase.

To support a polystyrene onto the upper rim of calix[4]arene (phenolic-0- of calix[4larene) and 25,26,27tribenzoyloxy-28-hydroxy, calixt4larene was treated with chloromethylated polystyrene in the presence of K2CO3 (Scheme 7). Polymeric calix[4larene (3a) thus obtained was hydrolyzed inthe benzoyl groups prior to use for the extraction process. Selective extraction experiments were then performed to see transference of some transition elements (Cu2+,Ni2+,CO’+, and Fe3+)from the aqueous phase to the organic phase by the synthesized polymeric calixarenes. Phase-transfer studies in water-chloroform confirmed that polymer Zb and 3b were Fe3+ ion-selective as was its monomer (1). Extraction of Fe3+ cation with 2b and 3b was observed to be maximum at pH 5.4. Only trace amounts of other metal cations such as Cu2+, Ni2+,and CO’+ were transferred from the aqueous to the organic phase (Table 3). Furthermore, the extracted quantities of these cations remained unaffected with increasing pH. The effect of pH on the extraction of 3b was lower and 56% extraction was accomplishedeven at pH 2.2. The extraction experiments were also performed with calix[4]arene(1); the ratio was 8.4%at pH 2.2.The polymeric calix[4]areneswere selective to extract Fe3 from an aqueous solution, which contained Cu2,Ni2 , CO’+, and Fe3+ cations, and it was observed that the +

+

l

J

2

3

+

346

Yilmaz

&&! OR

S

4a:COCH3 4b: H

(1,3alternate)

4

polymer supportedonto the lower rim calix[4]areneis a more efficient carrierfor Fe3+in the extractionprocess. The results of the extraction of metal cations are summarized in Table 3. The 4b showed the selectivity toward Fe3+ cation and cannot act as an extractant for Fe3+ at pH 2.2. The extraction of Fe3+ cation with 4b was increased with increasing pH, but the Fe3+ cation ratios at all pH decrease than that of 2b and 3b. The results suggested that the conformation of 4b and the C=O groups in the polymer affected the extraction of Fe3+ in this process. Based on the preceding results, we have observed that the extraction process with the polystyrene sup-

ported onto the lower rim calix[4larene (3b)can act cooperatively, whereas those induced onto the upper rim (i.e., open p-position side) act rather independently. In another study in this field, Deligoz et al. [SO], synthesizedapolymericcalixarenesbycombining 25,26,27-tribenzoyloxycalix[4larenewith the oligomer1 in the presence of NaH. Based on the chlorine analysis of this product, it was observed that compound 2 did not attach to each consecutive (CH*-Cl) groups in a regular array. The polymeric calixarene(2) thus obtained was hydrolyzed with ethanolic NaOH solution to remove the benzoyl groups prior to use for the extraction process.

Table 3 Extraction of Metal Cations with Ligands*

Metal cations extracted (%) Ligand 1 1 1 1

Ni2

+

+

8.4 56.0 57.5 5.4

90.0

9.6

l, and the break in the chain middle becomes more probable (denominator increases). At uniform length of nonassociated site it is more profitable, thermodynamically, for it to dispose at the end (number of states and, consequently, entropy will be higher for uniform length of nonassociated sites if the site is fixed by one end and not by both). This case promotes the reaction proceeding at chain ends at low temperatures.

Degradation Reaction Thermodynamic the to Opportunity Due

We will consider this question in more detail in the following text. Considering the influence of various parameters on T, it is necessary to take into account that a, nph, a‘, nLh depend on n, and on each other. Parameters a and a’ increase with the chain length at constant nph and nbh and decrease with nph and growth at constant n, value. As the change of a lies in relatively narrow ranges compared withnph, in most cases, the predominant influence on the degradation process is the amount and distribution of intermolecular bonds. With the chain length growth, the difference (the distribution width) between the maximumpossible amount of the intermolecular bonds and the minimal amount of those necessary for the chain retaining on the surface increases. It follows from this fact, that the probability of the equilibrium formation, as well as of stationary states at adsorption, occurs in dependence on storage conditions. Apparently, different amounts and distributions of intermolecular bonds is one of the main causes of stationary state formation. If nph - n t is (for example, molecule was fixed by both ends before the process, and then each fragment is fixed byone end), then a’is greater than a (flexibility of the chain, fixed by both ends, is usually lower than that of chains fixed byone end). In thiscase the denominator increases and T decreases with n, growth. Although it isclear that a grows with n,, the molecule can beunstable at any definite temperature at definite length. If the chain is stretched sufficiently by the surface (intermolecular bonds are sufficiently removed from each other), then we can accept for the first approximation a -+ 0, and a’+ 1. Then: A Hx T= AS, + R(n, - 3) At definite values of T, A H, and A S, from the equation (28), the length of nonassociated parts is determined at which degradation reaction is possible. At definite AH, and AS,, Twill be thermodynamically lower than the temperature at which the break of chemical bonds in low-molecular compound proceeds. The value of a at definite n, will be defined by the distance between engagement points(there exists optimal valueof this distance, at which a possesses maximum value).The value of Twill depend on engagement places, and apparently, on the disposition anddistribution of surface active centers. In the limit, A H, - TAS, = RTn, where n, for the first approximation equals the length of nonassociated sites of the chain. In thiscase it isnecessary to take into account that if T > Tb(Tb is the boiling temperature), molecules will evaporate, and if T < Tb, molecules will be decomposed until they form the chain fragments, being ableto evaporate. Analyzing degradation it is necessary to take into account that at temperature increase the value of free energy changeof bond break decreases, and the difference of EjGdefli and E:iGdefl,iincreases.

Initiation

369

The first term may increase in consequence of the growth of a’and T, and the second one is a consequence of temperature growth only. Considering the process proceeding onthe surface, the process of chemical degradation itself should be taken into account as well as the break of intermolecular bonds. However, at break of all intermolecular bonds there exists a limited lengthof thermodynamically stable chain onthe surface, which isstipulated by the impossibility of a large molecule evaporating. Thermodynamic instability is definedby the fact that if molecule occurs on the surface, all the states cannot be realized in it, as would occur in gas.The transition of molecules into the gas phasedepends on the attraction force. A limited size of thermodynamically stable chain willchange in dependence on thisattraction. In the absence of the attraction, molecules are able to transit into the gas, and their stability will be the same even for infinitely large molecules as for low-molecularcompounds. This can beeasily shown by modeling isolated molecules. Ifa molecule is fixed by two points on the surface, it represents something like cyclic compositions. Inthis case, system stability depends on fixing places of the chain itself onto the surface (Fig. 20). The disposition will be mostly stable when engagement places are situated nearby and closer to the chain end. Then the system may be considered as fixed by single end. Free energy of formulation will decrease according to translocation of fixation places because longer ends possess high a value (for infinitely long chains a does not practically depend on their length). Atthe decrease of chain ends, i.e., broadening engagementplaces, free energy of the initial system formation will decrease and become minimal when both ends are fixed, disposed from each other at the distance equal to contour chain length. In this state, a will be close to zero, and a’ will be sufficiently higher (in the limit a’+ 1). As twoends are formed at the chain

Figure 20 The dependence of deflection entropy related to monomeric unit (bond) Sdefl/n on the distance of fixed points. The molecule is fixed by two points on the surface.

370

Figure 21 Scheme of the probable breakof the molecule. Moleculeisfixedbytwo points on the surface. Arrows show the break direction. a = in the middle of the cycle; b = on the chain-surface bond.

break and their free energy of deflection depends on length, then the break will occur at the middle of the “cycle” according to the highest probability (Fig. 21). If changes of free energy of chain-surface bond formation are lower than that of chemical bond formation, then with regard to higher free energy of longer sites fixed by one end, chain-surface bonds will break with higher probability.This phenomenon must beobserved at the association withthe formation of physical bonds. If a molecule is fixed on the surface by three and more active centers, and if the energy of chain-surface bond is higher than bond energy in the macromolecule (chemical chain-surfacebond), degradation maybe considered for different cycles separately. In this case the break at the middle of the cycle will be morethermodynamically probable. If the bond energy in the chain is higher thanthat of chain-surface bond, then the reaction initiationproceeding by chain-surface bondismore probable. In this case, reactions forming longernonassociated ends, are preferable, i.e., the break of the “end” chain-surface bonds is more thermodynamically profitable (Fig. 22). However, in the case of probable formation of large cycle (Fig. 23)the break of chain-surface bond will be thermodynamically profitable for “internal” bonds, too, in consequence of high benefit of free

Figure 22 Scheme of the adsorbed molecule break. Arrows show direction of break: a = break of chain; b = detaching of the end of the chain-surface bond.

Vainstein and Zaikov

Figure 23 Scheme of thedetaching of a chain-surface bond with increasing length between two fixed points.

energy by means of the increase of the length of nonassociated site in the cycle, comparing with the increase of the length of the end site. For example, if the length of the end site increases by one unit and the large cycle is formed at the break of the middle bond, the reaction of cycle formation is thermodynamically preferable. The place of reaction proceedingdepends mainly onentropic components of free energy. Both reaction places are equally probable if free energy change is the same for anequalnumber of breaks of variously disposed chain-surface bonds. In this case, equal probability of the break may be observed at the reactions of cycle break and at the growth of the chain end, and at the reactions of cycle size increase. If lengths increase equally (for example, such a case may be displayed at equal probability of bond formation at any place of the chain), the increase of nonassociated end length is always preferable. That is why chain ends are less associated than its middle at adsorption, macromolecule association, and other similar processes. Calculation of combinatorial free energy change should be performed in the presence of relatively weak(physical, preferably) chain-surface bonds. The influence of this component change on the process is similar to that of excessive combinatorial entropy of different size molecules in solution. However, absolute value of excessive combinatorial entropy on the surface is usually lower than for the solution because the number of various dispositions of molecules in the solution is calculated for three dimensions, and on the surface for two dimensions. Solvent influence on the change of process free energy is displayed by practically all the terms, except (AGS,,, AGcom,s).Thus, estimating (G& - G,,,& it is necessary to take into account combinatorial analysis, too, as a consequence of the surrounding of nonassociated sites of chain molecules and contact with nonassociated molecules. As previously mentioned, the solvent makes the chain more rigid in most cases. That is whyinmany cases the (zjGdefl,j - ziGdefl,i) value can decrease. Moreover, the solvent changes local free energy of chain-surface bondformation, because at its break there may occur the gain of free energy of surface formation. In the presence of a solvent in the first approximation, it may be taken that all local places of the surface able

to formintermolecularbonds are occupiedby these bonds or by solvent molecules.That is why (G[d - Gid) will change weakly in comparison with their value in solvent absence, if the disposition of molecules on the surface is considered only. The (G& - Gid) will also depend on the system volume, as well as on thesurface size, if general combinatorial analysis of the system is calculated with regardto solvent molecules disposedon the surface. However, inthe last case it isnecessary to take into account the transfer of initial and degradated polymer moleculesas well as solvent ones to thevolume (into the solution). Moreover, it is necessary to introduce components offree energy change connected with the change of local interaction of the molecules surrounding the chain witheach other, and local interaction of associated and nonassociated parts of chains withthe solvent. The influence ofthese components is similarto their influence on molecule degradation inthe solution. Only (GL0m.S - Gcom+,)will be practically constant because this term depends just on the structure of chain and surface. The exception is in the case of changing local conformation, as well as the whole chain. The chapter did not consider the contribution of conformation variation into the change of the process of free energy. In general, case values of limited parameters at which there appears thermodynamic opportunity of reaction initiation decrease at the formation of thermodynamically unstable conformers. Thus, as it follows from the consideration of this text thermodynamic opportunity of reaction initiation occurs at the change of conditions. The performance of this reaction is defined by kinetic opportunities of the process. As the conditions of production usually differ sufficiently from those of storage and exploitation of polymers, then as a consequence of all the abovementioned, there occur the opportunities of degradation. This chapter estimated probabilities of degradation initiation for comparatively simple processes. In general, a similar approachcan be applied to thechange of reaction mechanisms, which can prevail in various conditions of storage and usage of the sample [27-301. REFERENCES N. M. Emanuel andA. L. Buchahenko,Chemical Physics of Polymer Degradationand Stabilization, VNU Science Press, p. 340 (1987). G . E. Zaikov, A. L. Iordanskii, and V. S. Markin, Diffusion of Electrolytes in Polymers, VSP Science Publisher Utrecht, p. 322 (1988). Yu. A. Slyapnikov, S. G. Kiryushkin, and A. P. Mar’in, Antioxidative Stabilization of Polymers, Ellis Horwood, Chichester (W. Sussex), p. 246 (1994). E. F. Vainstein, The Role of Intermolecular Interaction in Polymer Degradation Processes, Polymer Yearbook, (R. A. Pethrick and G. E.Zaikov, eds.) Gordon and Breach, London, vol. 9, pp. 69-79 (1993).

5. A. Popov,N.Rapoport,andG.Zaikov, Oxidation of Stressed Polymers, Gordon and Breach, London,p. 336 (1991). 6. K. S. Minsker, S. V. Kolesov, andG . E. Zaikov, Degradation and Stabilization ofPolymers on the Base of Vinylchloride, Pergamon Press, Oxford,p. 526 (1988). 7. E. F. Vainstein, The Limiting Sizes of the Large Molecules, Polymer Yearbook,(R. A. Pethrick, andG . E. Zaikov,eds.)GordonandBreach,London,vol. 8, pp. 85-113 (1992). 8. Yu. V. Moiseev and G. E. Zaikov, Chemical Resistance ofPolymers in Reactive Media, Plenum Press,New York, p. 586 (1987). 9. R. M. Aseeva and G. E. Zaikov, Combustion of Polymeric Materials, Karl Hanser Verlag, Miinchen, p. 390 (1986). ‘10. N. Grassie, and G.Scott, Polymer Degradation and Stabilization, CambridgeUniversityPress,Cambridge, p. 222 (1985). 11. E. F. Vainstein, A. A. Sokolovskii, andA. S. Kuzminskii, Kinetics of the Changing Products Molecular-Mass Distribution in Thermodegradation of Associated Polymers, Polymer Yearbook, (R. A. Pethrick and G . E. Zaikov, eds.) Gordon and Breach, London, vol. 9, pp. 79-101 (1993). 12. E. F. Vainstein, G. M. Gambarov, and S. G.Entelis, Doklady Akademii Nauk SSSR. 113(2): 375-378 (1974). 13. E. F. Vainstein, ThesisDr. of Sci., Instituteof Chemical Physics USSR Academy of Sciences, Moscow (1981). 14. L. D. Landau, and E.M. Lifshitz, Statistical Physics (in Russian), Nauka, Moscow(1964). 15. P. J. Flory, J . Chem. Phys., 10(1), 51 (1942). 16. H. L. Huggins, Ann. N.-Y. Academy of Sciences, 43 (Pt. 1): 1(1942). 17. E. M: Guggenheim, Mixtures, Clarendon Press, Oxford, p. 272 (1952). 18. A. J. Staverman,Rec. Trav. Chim. Phys.-Bas., 1 (1950). 19. I. Prigogine, The Molecular Theory of Solution, NorthHolland Publishing Co., Amsterdam, p. 448 (1957). 20. A. D. Pomogailo, ImmobilizedPolymeric-MetalComplex Catalysts (in Russian), Khimiya, Moscow (1986). 21. I. Oshima Takumi, and Nagai Toshikazu,J . Org. Chem., 56(2): 673 (1991). L. Knunyants, 22. Chemical Encyclopedian Dictionary, ed.) Sovetskaya Encyclopedia, Moscow, p. 791 (1983). 23. E. F. Vainstein and G . E. ZaikovProcesses with Oligomeres, Polymer Yearbook, (R. A. Pethrick and G. E. Zaikov, eds.) Gordon and Breach,London, 10: 231-235 (1993). 24. E. F. Vainstein, andA. A. Sokolovskii,11: 49-73 (1994). 25. 0.F. Shlensky, A. A. Matyukhin, and E. F. Vainstein, J . Thermal. Anal. 34: 645 (1988). 26. A.A. Sokolovskii, E. F. Vainstein,K. M. Gubeladze, and A. S. Kuzminskii, Vysokomolekulyarnye Soedineniya, 30B (1): 244-248 (1988). 27. G . E. Zaikov,Russian ChemicalReview, 62(6):603-623 (1993). 28. G . E. Zaikov, Intern. J . Polym. Mater. 24, (14): 1-19 (1994). 29. V. V. Kharitonov,B. L. Psikha, and G. E. Zaikov, Intern. J . Polym. Mater. 26 ( 3 - 4 ) : 121-176 (1994). 30. G. E. Zaikov, 208th American Chemical Society National

(I.

Meeting and Exposition Program, Washington DC, August 21-25, P. 125 (1994).


24 Polymerization by YIides A. K. Srivastava Harcourt Butler Technological Institute, Kanpur, India

1.

INTRODUCTION

Polymers in one form or another are basic constituents of every kind of living matter, whether plant or animal. It is only in the present century, however, that as a result of a wide range of scientific studies, their existence as a coherent group hascome to be recognizedand understood. Following this recognition and this understanding, the possibility has arisen of actually producing polymers by means of suitable chemical reactions. Originally these synthetic products, regarded as substituents for existing natural polymers such as rubber or silk, led to the introduction of a vast range of entirely new compounds in the field of plastics, rubber, and fibers, many of which have properties different from those existing in natural materials. At present, the study of polymers, though it includes the original natural polymers, tends to be dominatedby these synthetic materials. It is from the industrial development of such polymers that the main stimulus to scientific research has arisen. Progress in the chemistry of macromolecules results not only fromthe introduction of an increasing numbers of novel monomers from known polymerization reactions, but also from the application anddevelopment of new initiators. Investigations devoted to the latter occupy an important place in the polymer field andare of great significance;as a result of this, many compounds such as peroxides, hydroperoxides, redox initiators, and organometallic compounds, including ylides, have been reported. Over the past two decades, ylide-polymer chemistry has been anarea of great interest to chemists due to its theoretical implications and industrial significance. The first ylide wasprepared and isolated by Michaelis and Gimborn [l] in 1894. This was an isolated event, and the first flurry of activity in the field of ylides oc373

curred approximately 50 years later, probablystimulated by the development of the Wittig synthesis of olefins. The term “ylide” was first introduced by Wittig [2] in 1944. It represents the ending of the name given to nitrogen organic derivatives, trimethyl ammonium methylide; “yl” is the ending of organic radicals (methyl, phenyl) and suggests a free valency, and “ide” implies anionic character (acetylide, phenacylide). The ylides may be definedas dipolar compounds in which a carbanion is covalently bonded to a positively charged heteroatom. They are represented by the following general formula:

,c-0

1

8

x

whereXmay be P,As, Sb, S, N, Bi and T1. According to this, an ylide structure is characterized by two adjacent charged atoms, the carbon atom having an orbital uninvolved in the covalent bond occupied by an electron pair.

A. Classification The ylideshavebeenclassifiedon the basis of the heteroatom covalently bonded to the carbanion. Accordingly,they can be differentiated into nitrogen ylide (Scheme 2), sulfur ylide Scheme 3, phosphorus ylide Scheme4, arsenic ylide Scheme 5 , antimony ylide (Scheme 6), bismuth ylide(Scheme7) and thallium ylide (Scheme 8).

374

Srivastava

0

\S

I

vide a comprehensive guide to the literature reported from 1966 to 1995, thereby indicating the advances made during the past three decades in the polymer-ylide arena for the synthesis of polymers andthe scope of ylides in the polymerization field.

-c\

0

"P-c, I 0

'AS-

0

@/

A.Nitrogen

O,

Ylide

The effectiveness of ylides in the field of polymer science was first described in 1966 by George et al. [l l] who felt that 3- and 4-(bromo acetyl) pyridines, which contain both the a-haloketone and the pyridine nucleus in a single molecule, could bequaternized to polymeric quaternary salts and finallyto polymeric ylidesSchemes 9 and 10 by treating these polymeric salts with a base.

0

C
128"C) [14,151. Although reinforcingfillers ZnO Addin genii. 1 eral are known to adversely affect the strength properAdd iii. stearic acid 1 ties of TPEs [16], there are reports of reinforcement of iv. Add 1/3 stearate zinc 3 ionic thermoplastic elastomers by carbon black [17-201. Add v. 1/2 filler 3 This chapter reports the results of studies on the Add vi. 1/3 stearate zinc 3 physical, dynamic mechanical, and rheological behavior Add vii. 1/2 filler 3 of zinc oxide neutralized m-EPDM, particularly in the Addviii. 1/3 stearate zinc 3 presence of stearic acid and zinc special outstearate, with Sheet ix. 1 reference to the effects of precipitated silica filler. 441

EP6

100

id

Datta and De

442

Table 1 Formulations of Mixes and Molding Conditions [33]

Mix Number Ingredient EP4 EP5

EP3

EPO EP2

EP1 ~

m-EPDM ZnO 10 10 Stearic Zinc stearate DCP Molding temperature, 170 ("C) (min) Molding time,

100

-

10

10 1

-

2.0 150 8

20

2060

Table 2 Formulation of Silica-Filled Mixes" [27]

Mix Number Ingredients

so

S1

S2

m-EPDM ZnO Stearic acid Zinc stearate Silica

100

100 100

100

10 1 30

10

10

10

1

1 30 20

1

0

" All mixes were molded at 120°C for 20 min.

-

-

120120 20

Mixes were compression molded to form 2-mm thick sheets at a pressure of 10 MPa at 120°Cfor 20 min. The hardness was determinedas per American Society for Testing and Material (ASTMD 2240, 1986) and expressed in Shore A. The stress-strain properties were measured using dumbbell-shaped specimens according to ASTM D 412 (1987) ina Zwick universal testing machine (UTM), model 1445, at a crosshead speed of 500 mdmin. Hysteresis studies weremade according to ASTM D 412 (1980) by stretching the dumbbell-shaped specimens upto an extension of 200%.Tear strength was measured in the Zwick UTM 1445 using a 90" nick cut crescent sample according to ASTM D 624-86. The tension set at 100%extension was determined as per ASTM D 412 (1987). Measurement of dynamicmechanical properties was carriedout under tension mode usinga viscoelastometer, (Rheovibron DDV-111-EP, M/s. Orientec Corp., Tokyo, Japan). Sample size was 3.5 cm X 6.5 mm X 2 mm. Testing was carried out at a low amplitude, 0.025 mm, over a temperature range of - 100°C to +2OO"C. Heating rate was 1"Clmin and frequency of oscillation was 3.5 Hz or 110 Hz. Rheological and processability behaviors were studied in a Monsanto processability tester (Mm), which is an automatic high-pressure capillaryviscometer. The entire barrel and capillary are electrically heated with a microprocessor-based temperature controller [141. The

30 30 10

~~~~~

S3

30

1 10

-

120 120

10 1

1 30

-

-

120 2020

20

100 10 1

40

-

capillary lengthto diameter ratio was 30: 1. Preheat time for each samplewas 5 min. Reprocessability studies were made by repeated extrusions, up to four cycles, of the samples through MPTat different shear rates.

111.

RESULTS AND DISCUSSION

A.

Effect of Zinc Stearate

Physical properties of the different systems are listed in Table 3. The m-EPDM-DCP system (mix EPO) was cured to thesame level as them-EPDM-ZnO (mix EP1) by choosing molding times so that the torque increase (i.e., the difference between the maximum torque and the minimum torque) is the same in both cases. As compared with the m-EPDM-DCP system, the mEPDM-ZnO system shows higher hardness, modulus, tensile strength, and tear resistance presumably due to the presence of ionic clusters, which will be discussed later. The higher elongation at break in the mEPDM-ZnO system is believed to be due to the occurrence of stress-inducedion exchange, which causes lowering of stress concentration, resulting in highelongation [21]. Hysteresis loss followsthe order EPO < EP1 < EP2 < EP3. Hysteresis of the ionic system (EP1) is higher than the DCP cured system (EM). Reinforcing fillers, in general, are known to cause an increase in hysteresis of rubbers. Accordingly, it is believedthat the ionic aggregates in the m-EPDM-ZnO system act not only as crosslink sites, but also as fillers providing reinforcement to the matrix [22]. The effect becomes more pronounced on addition of stearic acid and zinc.stearate. Incorporation of stearic acid facilitates formation of ionic aggregates. I t has been reported that at ambient conditions crystalline zincstearate acts as a reinforcing filler in metal oxidecrosslinked carboxylic rubbers [231. Results of measurements of physical properties at 70°C reveal that thereinforcing abilityof zinc stearate diminishes at higher test temperatures, presumably due to the onset of melting of zinc stearate and the consequent plasticization.

Maleated Thermoplastic on Elastomer Ionic Based

EPDM Rubber

443

Table 3 PhysicalProperties"[33]

Mix Number EP5

Properties EP4

EP3

EP2EPO

MPa)modulus300% (0.76) 580 (512)

Elongation (%) (MPa)strength Tensile

N/cm) strength Tear X l@ Hysteresis (J/m2) loss Permanent set at 100% extension (Shore Hardness A)

(%)

(0.80) 143 (130) 3.3 8 40

EP1 1.70 (1.80) 850

(422) 4.40 (2.24) 332 (202) 4.5 11 47

2.20 (1 .90) 620 (503) 5.80 (2.20) 370 (211) 7.0 10 53

2.60 (1.90) 880 (303) 7.15 (1.91) 381 (24) 9.0 11 56

2.80 (2.06) 800 (422) 8.54 (2.34) 400 (347) 10.0 13 58

- 63I

425

2.90 (1.80) 760 (431) 7.67 (2.42) 418 (233) 10.7 13 62

2.90 (1S O ) 710 (440)

6.34 (1.88) 422 (220) 13.9 15 64

Values in the parentheses are the results of tests done at 70°C.

Figure1shows the variationofdynamicstorage modulus (E') with temperature. As compared with the m-EPDM-DCP system, the m-EPDM-ZnO system shows a broad rubbery plateau, due to physical crosslinks arisingout of the ionic aggregates. Addition of zinc stearate causes a gradual increase inE' atambient temperatures, but asharp fall inE' ata temperature exceeding 100°C, indicatingonset of transition fromthe rubbery state to theviscous flow state. Itis also evident that the rate of fallof E' increases with an increase in zinc stearate loading.It isbelieved that at room temperature zinc stearate strengthens the ionic aggregates causing reinforcement of the matrix resulting in higher E'. But at a temperature approaching its melting point (i.e., 128"C), zinc stearate plasticizes the ionic aggregates, presumably by solvation or exchange reaction. In the absence of zinc stearate, plasticization cannot occur and, accordingly, the transition fromthe rubbery state to the viscous flow state is not observableunder the studied conditions of temperature range and strain levels. The absence of the transition results ainbroad rubbery plateau the in mEPDM-ZnO and m-EPDM-ZnO-stearic acid systems. Figure 2 shows typical plotstan8 of versus temperature, from there it is evident that the glass to rubber transitionin the m-EPDM-ZnO system occurs at a slightly higher temperature than the m-EPDM-DCP system. Furthermore, incorporation ofstearic acid and zinc stearate shifts the glass-rubber transition temperature (TB)of the m-EPDM-ZnO system to thehigher temperature side. The results are summarized in Table 4. The value of tan8 at TB(i.e., tans,,,) decreases with the incorporation of stearic acid and zinc stearate. This indicates that with the introduction of stearic acid and zinc stearate, thechainsbecomeconsiderablytightened, and, therefore, the backbone relaxation relatedto TBoc-

I lo6 I

-150

I

I

*l13

l

+: IO

TEHPERATURE [*C1

Figure 1 Plots of dynamic storage modulus against temperature at 3.5 Hz 1331. EPO (-), EP1 (-*-), EP2 (-..-),EP3(-...-)EP4(.....),EP5( ---),EP6(------).

444

1.o

Datta and De

I-

I

a01

-150

A

I

I

-63+l13

I

I +25

I

I

+200

TEMPERATURE 1°C 1

Figure 2 Plots showing variation oftans against temperature[33]. EPO (-), EPl (-*-), EPlD (-x-),EP2 EP3 EPS (---).

The high temperature relaxation peak is believed to be due to onset of motion of chain segments firmly held by the ionic aggregates called clusters [24,25]. Figure 3 is a model proposedfor the different types of ionic aggregates in ionomers and the region of restricted mobility in the chain segment causing a biphasic structure [25]. Restricted mobility region is separatedfrom the "free" mobility region by a skin layer. Formation of the ionic clusters is favored by higher content, ion provided other factors remain the same. That the high temperature relaxation is due to ionic clusters and not by any other relaxation of the base polymer is confirmed by the absence of the peak in the m-EPDM-DCP system. Thus, it can be concluded that in the present ionomer system, zinc stearate plays a dual role. First, below its melting pointit reinforces the matrix and strengthens the ionic aggregatesand, second, at a highertemperature it results in solvation of the ionic aggregates and plasticizes the system, thus, facilitating the transition from the rubbery state to the viscous flow state [23]. Whilestudyingmeltrheologyin the MPT, it was observed that the m-EPDM-ZnO and m-

(-a**-),

,)-m*-(

curs at a slightly higher temperature with consequent lowering of tans at T g . In the high-temperature region (i.e., above room temperature) a second relaxation occurs, which is, however,notprominent in the mEPDM-ZnO and m-EPDM-ZnO-stearic acid systems. The high temperature relaxation becomes prominent in the presence of zinc stearate. Characteristic features of the second transition is that the transition temperature, designated as Ti, shifts to a highertemperature side with the incorporation of zinc stearate and the tansvalue at Ti increases. At higher loading of zincstearate (i.e., 30 and 40 phr) tans at gets overshadowed by the rise in tan6 due to an early onset of melting of zinc stearate. Table 4 Results of Dynamic Mechanical Studies

at 3.5 Hz [33] Glass-Rubber Transition Mix No.

Tg("C)

tan6 at TB

EPO EP1 EP2 EP3 EP4 EPS EP6

- 44.0 -41.8 -40.0 -38.6 - 38.2 - 36.8 -37.0

1.273 0.094 1.192 1.084 0.103 0.112 0.944 0.121 0.934 0.128 0.837 0.139 0.820

a = absent.

Ionic Transition Ti ("C)

tan6 at Ti

a

-

38.0 40.9 46.3 51.7 52.6 54.8 Figure 3 Model of biphasic structure in ionomer.

Thermoplastic Elastomer Ionic Based

445 Rubber onEPDM Maleated

EP3

FRACTURE

loading. It is also apparent that the shear stress dependence of viscosity decreases with increasing zinc stearate loading, since higher loading of zinc stearate provides greater extent of plasticization and enhances the flowability of the system. At high-processing temperatures molten zinc stearate is believed to solvate the ionic aggregates (Fig. 5) and high-shearforce applied may facilitate exchange reactions. Due to the interchange reaction of the intermolecular carboxylate salts (both bridging and non-bridging types) with zinc stearate, the intermolecular metal carboxylate bonds can break and result in formation of a new structure with ionicgroups attached to the stearate moiety at one end. Since intermolecular salt bridges cease to exist in the postexchange reaction stage, the chains become free to flow. Since the system isprocessed as thermoplastics are processed, its reprocessability was studied under repeated cycles of extrusion in the MPT. Results of reprocessability studies are shown in Table 5. It is evident that after the first cycle viscosity increased slightly, which may be due to the orientation effect. In the subsequent cycles viscosity remained almost constant. Therefore, it is concluded that the zinc stearate plasticized zinc saltof m-EPDM is melt processablejust asthermoplastics. B.

I

25 LOG T ( kPa1

I

3.0

Figure 4 Plots showing dependenceof apparent viscosity (v)on apparent shear stress ( 7 ) [33].

EPDM-ZnO-stearic acid systems could notbe extruded even at 190°C. This is notunexpected since the material, in the absence of zinc stearate, shows no transition from the rubbery state to the viscous flow state (Fig. 1). In the presence of 10 phr of zinc stearate, the mEPDM-ZnO-stearic acid system could beextruded but melt fracture occurred at a lower temperature (150°C) at all shear rates. At 160°C and 170"C, however, the extrudates showed meltfracture only at high shear conditions. At 20 phr loadingof zinc stearate, melt fracture of the extrudate occurred at high shear conditions at 150"C, but at higher temperatures no melt fracture occurred and the extrusion was smooth under all shear conditions. At 30 and 40 phr loadings of zinc stearate, the extrudates were smooth under all shear conditions at all temperatures. 1.7 Figure 4 shows the plots of apparent viscosity ver1.7 sus shear stress at 170°C for different mixes.It is evident that the materials behave as pseudoplastic fluids, and the viscosity decreases withincreasingzinc stearate

Effect of Precipitated Silica Filler

The physical properties of the different mixes are summarizedinTable 6. As expected, both hardness and modulus increase with an increase in filler loading. Incorporation of silica causes a gradual increase in tensile strength up to 20 phr of the filler, when the strength is almost twice that of the unfilled system. At 30 phr of filler loading, the tensile strength comes down to the level of 10phr, but it is still 1.5 X the value of the unfilled system. The increase in elongation at break on filler incorporation is believed to be due to a slippage of the polymer chains over the silica surface [261. Because of the thermo-reversible nature of the ionic crosslinks, tensile strength dropped down at an elevated test temperature (i.e., 70"C), but the silica filled system registered higher strength than the unfilled system even at 70°C. Table 5 Results on Reprocessability Studies [33] No. of ShearStressApparentViscosity (KPa) Cycles (7) KPa-sec

1 2

3 4

232.6 253.3 259.1 263.2

1.5

1.8

300% Modulus of Extrudate (MW 2.69 2.74 2.87 2.74

Mix number, EP5; temperature, 170°C; shear rate, 159 S-

I.

Dana

446

De

Figure 5 Probable mechanism of shear-induced exchange reactions during melt flow process [33]. (I) Interaction of zinc stearate, (RCOO)* Zn, with ionic aggregates before melt flow. (11) Exchange reactions during melt flow.

Table 6 PhysicalProperties" [27]

Mix Number

so (Shore Hardness69 300%(MPa) modulus

A)67

S1 S3

62 2.90 3.64 3.26 (1.82) (2.35) (2.58) Tensile(MPa) strength 7.60 14.35 10.30 (2.42) (3.59) (3.10) Elongation break at (%) 760 13401370 (431) (530) (400) (N/cm) strength Tear 418 436 (233) (296) (256) Hysteresis work, (J/m2) x 103 80.1 112.0 98.0 Tension set at 100% extension (%) 22222017

S2

65

" Values in parentheses are the resultsof tests done at 70°C.

4.49 (3.13) 10.43 (4.00) 1000

449

(440)

424 (283) 110.2

Maleated EPDM Rubber

Thermoplastic on Elastomer Based Ionic

447

Typical plots of the storage modulus versus temperature and mechanicalloss (tans) versus temperature are shown in Fig.6. It is clear that, apart from the glass-rubber transition (TB)occurring around -37"C, there occurred the ionic transition (Ti), at high temperature (>50"C), which, inthe presence of silica, becomes prominent and is shiftedto the higher temperature side. The high temperature transition is believedto be due to relaxation of the restricted chain segments arisingout of the ionic aggregates called multiplets and clusters, as depicted in Fig. 3. It is shown that the skin layer separates the restricted mobility region fromthe free mobility region, which constitutes the bulk of the chain segments. Results of dynamic mechanicalstudies are summarized in Table 7.

D

l m

z

a

l-

01

."-100

I

I

I

I

I

200

100

0

TEMPERATURE ['C

I

Figure 6 Typical plots from dynamic mechanical thermal analysis showing storage modulus andtans variation with temperature [27]. SO (-), S2 (------).

101 I

Both tear resistance and hysteresis increase on incorporation of silica, but the effect isless pronounced as compared to the stress-strain properties. Tension set of the ZnO-neutralized m-EPDM system is low (around 20%) and incorporation of filler causes only a marginal increase in set due to chain slippageover the filler surface, as previously discussed. Measurement of physical properties reveal that there occurs an interaction between the filler' surface and the polymer. Results of dynamic mechanical studies, subsequentlydiscussed, support the conclusions derived from other physical properties.

F.

W

Table 7 Results of Dynamic Mechanical Studies [27] Mix Number

T,, "C

tans at T,

x,"C

tans at Ti

so

-36.8 -37.0 -36.4 -35.5

0.837 0.814 0.715 0.692

+52.6 +60.6 +62.9 +72.8

0.128 0.139 0.148 0.166

S1 S2 S3

0.70

am am

a06 1 P1

004

008

0.10

Figure 7 (a) Variation of EJ/E; with volume fraction [8] of the filler at room temperature [27]. (b) Variation of tans,/ tans, with volume fraction [8] of filler [27]. T,, (---m---) Ti (

0 v

1.

'

448

Datta and De Free mobility region

Figure 8 (a) Schematic diagram showing distribution of fillers in different parts of anionic elastomer [27]. (b) Proposed structural model showing interaction of silanol groups on silica surface with carboxylate groups [27].

The variation in roomtemperature storage modulus The results could be fitted into the following equation [27]:

on filler incorporation is shown in Fig. 7.

where E j is the storage modulus of the silica filled system and E; refers to themodulus of the gum or unfilled system and [e] is the volume fraction of silica. This is very similar to the relationship proposed by Smallwood [28] in the case of diene rubbers. The greater slope of the plot in Fig.7, ascompared with conventionalrubber systems, is ascribed to the strong interaction between

the polymer chains containing ionic aggregates and the active sites on the filler surface. Figure 7 also showsthe plots of (tanS)f/(tanS), versus volume fraction of the filler ( 0 ) at Tg and Ti. Here “f” stands for the silica filled system and “g” denotes the gum or unfilled system. The results could be fitted into the following relations [27]: At T g ,

Equation (2) is similar to that obtained in the case of conventional rubber systems anddepictsweakrub-

ubber EPDM Maleated Thermoplastic on Elastomer Ionic Based

449

Table 9 Results of Reprocessability Studies [271

Shear Stress Cycle 3.82 3.98

293.27 296.40 297.29

3.91

1 2 3 4

300% Modulus of Extrudate (MPa)

285.71 4.00

Mix no., S2; temperature, 170°C; shear rate, 121 S".

T IkPal

Figure 9 Variation of apparent viscosity [77] with shear stress [TI [27].

ber-filler interaction [29,30], involving the backbone chains of the polymer. At Ti,

At Ti the slope in the plot of Eq. (3) is positive, thereby 'inferringthat an increase in filler loadingstrengthens the high temperature relaxation, presumably by increasing the ionic cluster-induced chain rigidity. While the rubber-filler interaction involving the nonionicpolymer backbone is of weak van der Waals type, the same due to ionic aggregates can be of both hydrogen bonded type involving the polar sites on the filler surface [l31 and van der Waals type, as depicted in Fig. 8. Mondal et al. [3 l] made a similar observation in the case of zinc oxide neutralized carboxylated nitrile rubber filled with precipitated silica filler. Figure 9 shows the variation of apparent viscosity with apparent shear stress. It is evident that the mixes are pseudoplastic in nature. Furthermore, as expected, viscosity increases with an increasing filler loading. Activation energy(A E ) for viscous flow was calculated according to the following equation [32]. q = A e AEIRT

This indicates that the occurrence of the high temperature relaxation of the restricted mobility zones of the chain segmentsis facilitated by the filler, thereby implying strong bonding betweenthe ionic aggregates andthe polar sites of the filler. From the above results, it can be inferredthat the rubber-filler interaction, in the case of silica filled ionicrubber, is of two types: (1) the interaction between the filler particles and the nonionic segments of the polymer backbone, which is similarto the interaction involving diene rubbers and reinforcing fillers, as manifested in the lowering of tan8 at Tg; and (b) the interaction between the ionic groupsof the polymer and the polar sites (or silanol groups)present on the filler surface, which is manifestedin an increase in tans atTi. Table 8 Activation Energy (AE) for Viscous Flow (kJ/mol) [27]

Mix Number (sec") rate Shear 61 305 610

SO

S1

S2

S3

16.0 12.2 9.36

28.1 20.5 15.6

31.8 24.3 16.9

39.5 28.2 18.6

(4)

Here 7 is the apparent viscosity at temperature T, R is the universal gas constant, and A is an empirical constant, called frequency factor for melt flow.The activation energy valuesfor different systems and at different shear rates are summarized in Table 8. It is evident that activation energyfor flow increases with filler loading, but it decreases with an increase in shear rate. At high-processing temperature, zinc stearate plasticizes the matrix andsolvates the ionic clusters [28,29], thus facilitating the polymer flow.In the silica filledsystems, however, the flow is hindereddue to thepresence of silica particles anchoring the chain segments both in the regions of free mobilityand restricted mobility, thereby increasingthe free energy of activation. Furthermore, under high-shear conditions, it is likely that the loosely held silica aggregatesare broken down, causing a reduction in the apparent activation energy of flow. Since the system exhibits thermoplastic behavior, its reprocessability was studied in the MPT and the results are given in Table 9. Viscosity duringrepeated extrusions and the modulus of the corresponding extrudates remain constant up to four cycles of extrusion, thereby confirming the thermoplastic elastomeric nature of the silica filled m-EPDM-ZnO-stearic acid-zincstearate composition.

Datta and De

450

W . CONCLUSIONS

Zinc salt of maleated EPDM rubber in the presence of stearic acid and zincstearate behaves as a thermoplastic elastomer, which can be reinforcedby the incorporation of precipitated silica filler. It is believed that besides the dispersive type of forces operative in the interaction between the backbone chains and the filler particles, the ionic domainsin the polymer interact strongly with the polar sites on the filler surface through formation of hydrogen bonded structures. REFERENCES 1. R.A.

Weiss, J. A. Fitzerald, and D. Kim, Macromolecules, 24: 1071 (1991). 2. M. Pineri, C. Mayer, A.M. Levelut, and M. Lambert, J . Polym. Sci.. Polym. Phys., 12 : 115 (1974). 3. R. Jerome, J. Horrion. R. Fayt, and Ph. Tayssie, Macromolecules, 17: 2447 (1984). 4. U. K. Mondal, D. K. Tripathy, and S. K. De, Polymer, 34: 3832 (1993).

K. Sato, Rubber Chem. Technol., 56: 1942(1984). R. A. Weiss, W. J. MacKnight, R.D. Lundberg, K.A. Mauritz, C. Thies, andD. A. Brant, “Coulombic Interactions in Macromolecular Systems,” A.C. S. Symposium Series, 302: 1 (1986). 7. P. K. Aganval and R. D. Lundberg, Macromolecules, 17:

5. 6.

1918 (1984).

8. P. K. Agarkal and R. D. Lundberg, Macromolecules, 17: 1928 (1984). . . 9. A. U. Paeglis and F. X. Oshea, Rubber Chem. Technol., 60: 223 (1988). 10. A. J. Oostenbrink and R. J. Gaymans, Polymer, 33: 3086 (1992). 11. D. R. Paul and J. W. Barlow, Polymer, 33: 268 (1992). 12. R. Greco, M. Malineonico, E. Martuscelli, G. Rogosta, and G. Scarizi. Polymer, 28: 1185 (1987). 13. E. N. Kresge,“Ionic Bondingin ElastomericNet-

14. 15. 16. 17.

works,” 18th Canadian High Polymer Forum, Hamilton, Ontario (1978). U. K. Mondal,D.K. Tripathy,and S. K.De, Polym. Engg. Sci. 36:283 (1996). I. Duvedevani, P.K.Agarwal, and R.D. Lundberg, Polym. Eng. Sci., 22: 500 (1982). E. J. @an, Handbook of Fillers and Reinforcements for Plastics, (H. S. Katz and J. V. Milewski, ed.), Van Nostrand Reinhold Co., New York (1978). U. K. Mondal, D.K. Tripathy, and S. K. De, Plastics Rubber and Composites-Processing and Applications,

24: 19 (1995). 18. K. Sato, Rubber Chem. Technol., 56: 942(1983). 19. A. U.Paeglis, and F. X. 0.Sheha, Rubber Chem. Technol., 60: 228 (1988). 20. T. Kurian, P. P. De, D. Khastgir, D. K. Tripathy, S . K. De, and D. G . Peiffer, Polymer 36: 3875 (1995). 21. H. P. Brown and C. F. Gibbs, Rubber Chem. Technol., 28: 937 (1955). 22. B. HirdandA. Eisenberg, Macromolecules, 25: 6466 (1992). 23. M. R. Tant and G. L. Wilkes, J . Macromolecular Sci. Rev., Pt. C, 28: 1 (1988). 24. A. Eisenberg, B. Hird,and R. B. Moore, Macromolecules, 23 : 4098 (1990). 25. W. J.MacKnight, W. P. Taggert,and R. S. Stein, J . Polym. Sci., Polym. Symp., 45; 113 (1974). 26. S. K. Chakraborty,and S. K. De, Polymer, 24: 1055 (1983). 27. S. Dutta, A. K. Bhattacharya, S. K. De, E. G. Kontos, and J. M. Wefer, Polymer 37: 2581 (19%). 28. H. M. Smallwood, J . Appl. Phys., 15: 758(1944). 29. L. E. Nielsen, J . Polym. Sci., Polym. Phys. Ed.,17: 1897 (1979). 30. D. Rov. A. K. Bhowmick, and S. K. De, J . Appl. Polym. Sci., i 9:263 (1993). 31. U. K. Mondal, D. K. Tripathy, and S. K. De, J . Appl. Polym. Sci., 55: 1185 (1995). 32. A. A. Tager, Physical Properties of Polymers, Mir Publishers, Moscow (1978). 33. S. Dutta, S. K. De, E. G. Kontos, and J. M. Wefer, J . Appl. Polym. Sci. 61: 177 (1996).

Elastomeric Compound-The Importance of Consistency in Economic Component Manufacture Robert Bond Chicago Rawhide (Americas), Elgin, Illinois

1.

INTRODUCTION

Figure 1 shows a typical rubber covered radial lip shaft seal. It is made up of three basic components, the metal insert, the garter spring, and the elastomeric lip. Following isa description of the role of each component.

Today with the everincreasing global market place and the demands for customer satisfaction, the end user is placing more and more emphasis on product quality, reliability, and value. The result is that companies that produce components for the major original equipment manufacturers (OEMs), particularly the automotive industry, are being required to: improvequality, improve reliability, increase product life, improvecost effectiveness, and reduce response time. To this end, manufacturing throughput time has to be minimized.This allows the flexibility requiredto keep the response time to customers down to a minimum. The introduction and use of electronic data interchange (EDI) for changing order schedules, in some cases on a daily basis, make itessential for component manufacturers to have responsive and flexibleprocesses. One of the major contributingfactors that enable these objectives to be achieved is consistency, whether it be the incoming raw materialor the processes used in the manufacturing process. To get a more in-depth understanding of the issues relatingto raw materials and processes, the manufacture of radial lipshaft seals will be usedas thecomponent under consideration.

Metal insert: The metal insert acts as the stiffening element of the seal, which is essential when the seal is fitted into the bore or housing. Elastomeric lip: The elastomeric lip forms half the dynamic sealinginterface, the other half being the shaft. In the case of the rubber covered seal, the elastomeric compound also acts as half the static seal between the borehousing and the seal; the bore itself being the other half of the static seal. Garter spring: The garter spring generates the majority of the radial load the elastomeric lip exerts on the shaft, which is essential when the lip has to follow the shaft when operating under dynamic conditions.

II. THE RADIAL LIP SHAFT SEAL The radial lip shaft seal is a dynamic sealing device. In general, its role is to seal fluids in and contaminants out when sealing againsta rotating, reciprocating, or oscillating shaft. The seal design can take many forms, e.g., with or without a garter spring; withor without dirt exclusion lips; metal case, rubber covered, or half m e t d half rubber.

Figure 1 Rubber covered lip type rotary shaft seals. 451

Bond

452

It can be seenthat the role of the radial lip seal and the tolerances to which it has to be manufactured make it an engineered product that is used in the engineering industry; a radial lipshaft seal is not a rubber component to use typical rubber tolerances that happen to be used by the engineering industry. 111.

MANUFACTURING PROCESS

The manufacture of elastomeric radial lip seals centers around the molding operation. Figure 2 [l] illustrates the manufacturing flow diagram for the two primary processes used: compression molding and injection molding. To minimize the cost of manufacture, one approach is to minimize the manufacturing “throughput” time, sometimes called “just in time” or “the Toyota lean manufacturing process.” The result will be minimum work-in-process (WIP) through the elimination of “queue time” and a significant reduction in the cost of quality when comparedto the moreconventional approaches to manufacturing used in recent years. However, the new approaches to manufacturing demandthree things fromthe raw materi-

als, purchased components, and the individualprocesses: high quality, consistency, reliability. Therefore, every effort must be placed onensuring the quality, consistency, and reliability of supply of the incomingrawmaterialsand components. Since this chapter is about the need for consistent, high-quality elastomeric compound, emphasis will be placed on: incomingrawmaterials,compoundmixing, extrusion (where appropriate), and molding. A. Incoming Raw Materials The requirement for consistent, high-quality incoming raw materials comes from the need for the elastomeric compounds to have consistent processing properties. Therefore, it is important to: 1. Understand the critical properties of the mixed compound that affect processability. It should be understood that the properties and acceptable operating values of those properties, which control the compound processability, are dependent on the processes under consideration, e.g., hotor cold feed extrusion, compression or injection molding.

1 1 1

1 Figure 2 Typical process flow diagrams: (a) compression molding, (b) injection molding.

scosity Mooney

453

Elastomeric Compound 2. Establish the contribution the individual ingreviscosity dients in the elastomeric compounds make to both the Mooney vulcanized compoundproperties and the uncured properties required for consistent processability. 3. Establish realistic, workable individual ingrediviscosity ent specifications. Mooney At the present time, the difficulty is correlating incoming raw ingredient properties with processability. Elastomeric compounds can be considered to be made up of four components: polymers, fillers, process aids, and other chemicals. During recent years a considerable amount of research has been undertaken to understand what in the makeup of a polymer affects the processability. In the late 1980s, the Rubber Manufacturers Association in the United States undertook a research project with the Department of Polymer Engineering at the University of Akron to evaluate the laboratory equipment available using specially made butadiene-acrylonitrile polymers withdifferent acrylonitrile levels, molecularweights, and molecular weightdistributions. The results from the studyconfirmed that, from the processingvariables viewpoint, the major factors are: frequency (shear rate), temperature (temperature), and deformation (strain). Therefore, if processability is to be measured on a regular basis, it would be extremely useful if a piece of equipmentwasavailable that could measure the dynamic properties under realistic operating conditions. Fortunately, one piece of test equipment has been developed, which is commercially available, the RPA 2000 (Monsanto Co.), which may meet the requirements. A considerable number of investigationshavebeen reported on the RPA 2000 [2],that support the view that it maymeet the requirements of an instrument that measures both polymer and compound processability. The work to date identifies differencesin polymers and compounds. However, it is important to relate those differences to processing characteristics in the manufacturing environment. An alternate approach was taken by a shaft seal manufacturer anda polymer producer working together [2].The project involved makinga series of butadieneacrylonitrile polymers of known acrylonitrile contents with various molecular weights and molecular weight distributions then making seals under normal manufacturing conditions. There were four polymers evaluated:

A. Bound acrylonitrile Average molecular weight Intrinsic viscosity Compound Mooney B. Bound acrylonitrile weight molecular Average Intrinsic Compound C. Bound acrylonitrile

379,700 1.44 56 34% 306,200 1.13 37 Note: (1) Polymer D was made up of blending A and C (50:50) in the latex stage. The result was a bimodal molecular weight distribution having an average value close to B; (2) All compounds used the same formulation.

weight molecular Average Intrinsic Compound D. Bound acrylonitrile weight molecular Average Intrinsic Compound

The performance of the differentpolymerswas based onthe molding scrap levels obtainedfor two categories of scrap: Excess flash-Indicates the compound flow was excessive during the mold closing process (shelf loaded compression molding) [31. Unfills-Indicates that the compound did not flow sufficientlyduring the mold closing process (shelf loaded compression molding) [3]. The variable usedin the manufacturing process was the mold closure speed to investigate the effect of varying the shear rate (frequency). Figures 3, 4 and 5 illustrate the unfill scrap for the polymers A through D for mold closure times of 5 S, 8 S, and 11 S, respectively. Figure 6 shows the excess flash scrap for the same polymers with a mold closure time of 11 s. The results show that for butadiene-acrylonitrile polymers,molecular weight and molecular weightdistribution have a significant effect on the compound flow during the molding process: Inaddition, it shows that shear rate (frequency) plays an important role.

34% 275,500 1.05

28 34% 309,600 1.15 37 34%

"

A

B

C

D

Figure 3 Percent unfills by elastomer type for

closure.

5-Smold

Bond

r

I B

Figure 4 Percent unflls by elastomer type for 8-S mold

closure.

Unfortunately, at the present time there does not appear to be a detailed understandingof the relationship between the properties of the raw materialsor the ingredients making up an elastomericcompound andthe compound's processability.Therefore, the most widely used approach presently being adopted is to use the conventional characteristics with the standard specification limits issued by the supplier or special agreed upon limits with the supplier to accommodate specific requirements in the fabricators manufacturing process.

Figure 5 Percent unfills by elastomer type for 11-S mold

closure.

C

D

Figure 6 Percent excess flashby elastomer typefor 1l-S mold closure.

B. Compound Mixing

The previous section has discussed the importance of understanding the role the individual ingredients play in the compound processability. The next major issue is the mixing of the elastomeric compound. In any manufacturing operation that uses rubber (elastomeric compound) as one of the major components in the product being produced, the consistency of the mixed elastomeric compoundis critical not onlyfor product reliability and performance, but also for keeping the cost of quality down and enabling the manufacturing throughput time to be minimized.Therefore, the major requirements from the mixing operation are: flexibility, high-quality mixed compound, compound consistency (within batch and batchto batch), compound delivery, and cost effectiveness. There are two types of mixing processes available, continuous and batch. The nature of the component manufacturing business generally dictates that there is a need to mix a large variety of compounds. Therefore, the majority of component manufacturers use the batch mixing process to produce the different elastomeric compounds required. This process consists of four basic areas all of which are important for consistent mixing: ingredient weighing,batch mixing process, take-off system, and cooling and batch-offsystem.

l . Ingredient Weighing The systems used for ingredient weighingcan vary from a totally manual system to a totally automated system where the ingredients are untouched by human hands. The critical role of the ingredientweighing system,

455

Elastomeric Compound

whether it is manual or fully automated, is to ensure that the correct weight of the correct ingredient arrives at the mixer hopper charging door in the correct order at the right time. The least costly approach is to use a manual weighing system.However, this approach also has the highest risk for manual error, although witha very experienced mixing crew the problem can be minimized. The other extreme, a fully automated weighing system, is the most expensive but hasthe advantage that the possibility for human error is removed. This applies not only to the weighing process itself butalso ‘to the filling of the bulk containers, the critical ingredient containers, and the other chemical containers, including oils and greases if the incoming ingredients havethe appropriate bar code label attached. The bar code labels are read prior tofilling the container; if the bar code label corresponds to the correct ingredient thenthe container can be opened, if it is incorrect, the containers will not open. The system generally consideredto be the most appropriate is a combination of both approaches, a semiautomated computer prompted weighing system using the locked containerbar code labelling technique to minimize, hopefully eliminate, the possibility of getting the incorrect material in a given container. For consistent, high-quality compounds, both within batch and from batch to batch, it cannot be overstated that the correct weight (tight tolerances) of the correct ingredient must arrive at the mixer hopper charging door at the right time in the correct order.

2. Compound Mixing A cross-section of a typical batch mixer can be seen in Fig. 7 [4]. Using a simple analogy,the batch mixing process can be saidto be similarto mixing a fine fruit cake. The ingredients are placed in the mixing chamber then mixed together using two mixing rotors and finally discharged from the mixer. There are two fundamentalrotor designs in general use, the tangential rotor (Fig. 8a [51) and the intermeshing rotor (Fig. 8b[5]) designs. In recent years, a number of in-depth studies have been completed comparingthe mixing capability of the tangential rotor to the intermeshing design. Bond and Takada [5] in 1992reported that in both cases the latest (at that time) intermeshing and two-wing tangentialrotor mixers gavesuperior temperature control and mixing capability when compared to establish to the conventional 3D Banbury* mixer used the base line results. It also demonstrated the role the use of computer controls played in improvingthe mixed compound consistency, bothwithin batch and from batch to batch. Therefore, there are additional important factors that have to be taken into account when purchas-

* Banbury is a trademark of Farrell Corporation.

22” Diameter Welght Cyllnder Air-Operated

Floating Welght-

“Bedplate

Figure 7 Typicalbatchmixer.

ing a new mixer, e.g., equipment reliability, spare part availability, and technical service. 3. Take-OffSystem

When the compound is dropped out of the mixer, it is essential to remove the heat from the batch of rubber as quickly as possible; thisis particularly important when the curatives are in the compound. It is just as important to ensure that the batch is fully blended before it is finally cooled and stored prior to use. There are three systems in use today:

1. A twin-screw extruder fitted with a “stuffer box.” The twin-screw extruder continues to blend the compound as it passes through; although withthe dryer,

Figure8 Internal mixer rotor designs: (a) tangential rotor, (b) intermeshingrotor.

Bond

456

more difficult compound formulations can have problems. If the heat is to be removed efficiently,the extruder must have extremely good cooling capabilities. 2. A single drop mill sized appropriately for the mixer. With the single, two-rolldrop mill approach, the mill can be either beneath the mixer or alongside it and fed using a tub on railsbeneath the mixer. This enables the heat to be takenout of the compound fromthe mixer. If time is not a critical issue in the mixing process, the single drop mill approach also accommodates the blending required. 3. A two-mill take-off system, the mills sized appropriately for the mixer. The two two-rollmill take-off system is madeup of one mill beneath the mixer to take out the heat from the compound (the drop mill) and a second mill used for the final blendingprocess. This system is recommended when time and productivity are major issues.

5 . Optimized Mixing Line with a Computer Prompted Weighing System

To ensure that each batch of compound gets the same amount of blending (energy input), the blending mill is generally fitted with a stock blender to improve the blending consistency. To ensure that each batch of compound gets the same amount of blending, automated controls can beincorporated on the blending mill to control the following variables:

In the two typical shaft seal manufacturing processes shown in Fig. 2, the extrusion process is only required if the compression moldingprocess is beingused. There are basically two types of extruders in general use, the cold feed andthe hot feed. In both cases their role is to produce the preforms for use with compression molding. The critical issues with the preform are: the shape, the dimensions, and the weight.

Number of passes on the stock blender. Individual roll speeds. Individual roll temperature. “Nip” setting (gap between rolls).

Bond and Takada [5] showed how replacing a conventional, manual mixing line basedon a 3D Banbury mixer with a mixing line comprised of a computer prompted weighing system using bar code labels and readers, the latest 80-L Banbury mixer, a two two-roll mill take-off system, and conventional coolingandbatch-offimproved compound consistency. Figure 9 schematically illustrates the improved mixing line. Figures 10-14 illustrate the reductions obtained in hardness, 50% modulus, 100% modulus, elongation, and tensile strength, respectively, for a fluoroelastomer compoundproduced under normal production conditions. Table1shows the improvement obtained in detail. C. Extrusion

I . ColdFeedExtrusion Cold feed extruders use compounds in a strip form. The strip is generally made during the take-off process in mixing; a series of cutters are used on the blending mill

4. Cooling and Batch-Off System The conventional system in general use takes the compound fromthe blending mill, throughan antistick applicator onto a festoon system withair cooling, and finally through a “wig-wag”take-off system into storage baskets.

Figure 9 Flow diagram: improved mixing line.

Figure 10 Comparison of hardness -c 3 sigma variability ratings: fluoroelastomer compound.

457


OldMiikg Line

New M i k g Line

'

Figure 11 Comparison of 50% modulus k 3 sigmavariability ratings: fluoroelastomer compound.

to producethe required strip width, then passed through the cooling system and batched-off into either baskets or wound onto coils for an easier feed into the extruder. There are two types of cold feedextruders, the standard screw design type (Fig. 15) and one with a screw feed using a ram action to feed the extruder. In both

Figure 13 Comparisonof elongation +3 sigma variability ratings: fluoroelastomer compound.

cases the preforms are produced using a rotating blade system immediatelyafter the extruder die. If consistency is critical, it is important to remember that the preforms are only as consistent as the compound or rubber strip feeding the extruder. Because of the critical nature of the preform dimensions,extruder manufac-

1 g

c

Figure 12 Comparison of 100% modulus +3 sigma variability ratings: fluoroelastomer compound.

(60)

Figure 14 Comparison of tensile strength + 3 sigma variability ratings: fluoroelastomer compound.

Bond

458 Table 1 Comparison of Fluoroelastomer Compound

Properties Old mixing line f 3 sigma rating July-Sept.

Property Hardness (IRHD) 50% modulus 100% modulus Elongation Tensile strength

1989

New mixing line 2 3 sigma rating July-Sept. 1990

100 100 100 100 100

56.4 54.0 59.0 46.0 52.0

turers have developed machine controls that allow the compound to leave the extruder at a constant rate. However, the rate varies depending on the nature and type of compound being extruded. The need for this type of control system is to minimize any variation or inconsistency within the compound being extruded. After leavingthe extruder, the preforms are cooled. There are two types of cooling processes used. The choice depends on the compound type being extruded. 1. If the compound beingextruded is not attacked by water, for example, conventional butadieneacrylonitrile based compounds, the preforms are passed througha water cooling system. The water cooling system can be using either the spray technique or complete submersion. 2. If the compound can be attacked by water, for example polyacrylate and fluoroelastomer elastomer based compounds, the preforms are air cooled.

Figure 15 Coldfeed extruder.

2. Hot Feed Extrusion There are two types of hot feed extruders, one similar to the screw type cold feed extruder, except that the strip feeding the extruder is taken directly off the tworoll mill. The mill is used to further blend and heat the compound. It is then taken off the mill in strip form and fed directly into the extruder. The second.type of hot feed extruder is based on the action of a ram being used to feed the compound into the extruder die (Fig. 16). The compound is placed on a two-roll mill to further blend and heat rubber. It is then taken off the mill in “pig” form to fill the extruder chamber. In both cases, the preforms are produced using the rotating cutter system as described for thecold feed extruders. The preforms are also cooled the same way.

D. Molding

In the manufacture of shaft seals, the molding is the process where the compound andtreated metal stamping (phosphated and adhesive applied) are brought together to form the conventional lip type shaft seal. Therefore, it is criticalthat the compounds used are formulated to suit the molding process. There are three fundamental molding processes in use today: compression, injection, and injectiodtransfer. The choice of process is generally based on economic issues such as: volume of seals to be produced; elastomeric compound, cost and waste; and product design. There are four critical issues required for all the molding processes listed previously, they are [6]: Temperature distribution across the mold (top and bottom). Pressure distribution across the mold.

Figure 16 Hotfeed extruder.


Application of vacuum, whether directly through the mold or completely evacuating the chamber. Press controls. The variation intemperature across the top and bottom halves of the mold is a function of the press platen size, the flatness of the mold and platen surfaces at the two mold/platen interfaces, the mold construction, and the platen heating system. The larger the platen size, the more difficult it toismaintain a constant temperature across the platen. Therefore, one solution to the problem is to use a large numberof molding presses with a relatively small platen size, for example, four presses with platens 355 x 355 mm will have approximatelythe same production capacity as one press with a 710 x 710 mm platen.In reality, for high-volume production, large presses are more economical from the cost point of view and the floor space required. The two accepted methods usedfor heating platens are both electric: resistance heatingandinduction heating. Experience has shown that the mostefficient method for minimizing temperature variation across a larger platen is the use of multiinduction heating coils. With a three zone system on a 710 x 710 mm platen correctly balanced, it is possibleto maintain the temperature variation across the platen to within & 2.5”C. The pressure distribution across the mold is a function of both the mold surface and the platen flatness. The larger the mold size and the platen, the more difficult it isto maintain the high degree of flatness required. The mold and platen flatness is also critical in the area of heat transfer from the platen to the mold itself. During the shaft seal molding process, the elastomeric compound has relatively large distances to travel/ flow in the mold, particularly when using shelf loaded compression molding (Fig. 17a) and injection molding (Fig. 17b). History has shown that without the use of a vacuum, either applied directly through the mold or

Figure 17 (1)“Diaphragm”gatingsystemthroughprimary lip. (2) “Diaphragm” gating system through dirt lip. (3)“Point” gating system through seal outside diameter. (4) “Fan” gating system through seal outside diameter.

459

completelyevacuating the mold chamber, “air trapping” is a major problem. Therefore, the trend started in the United States of evacuating the mold/mold chamber has becomethe accepted way to mold liptype rotary shaft seals. Press control is critical; it isessential that the elastomeric compoundreaches the required cure state to optimize product performance yet remains in the press the shortest time periodto maximize productivity. To meet this objective, both compression and injection presses now use microprocessor controls, which enable variations in platen temperatures and compoundcure characteristics to be accommodated without sacrificing product performance or productivity. One compound formulation not is suitable for all the different molding techniques used. It was reported by Berens and Born[71 and Horve [8] that molded lip seals produced using “shelf” loaded preforms, as shown in Fig.17a,gavemore consistent product performance than the “cavity” loaded approach, also shown in Fig. 17a. However, for “trimmed” lip seals, both “shelf” and “cavity” loadedpreformsgavesimilarperformance.Whenconsideringinjectionmolding, various “gating” systems, as shown in Fig. 17b, have been successfully used for producing “molded” lip type rotary shaft seals. The reason being that the elastomeric compound formulation plays a major role in the compound flow and particle orientation. The same compound formulation is not suitable for all the different approaches to injecting the rubber into the mold. I. Compression Molding Compression presses are producedin various sizes throughout the world. In general, shaft seals are manufactured on presses with platensizes in the range of 350 x 350 mm to 715 X 715mm with clamp loads in the range of 1500 kN to 4000 kN. Figure 18 shows a typicalline of compression presses fitted with vacuum chambers. This type of press can be purchased with platen sizes in the range 355 x 355 mm to 710 X 710 mm. The choice of press size is dependent on: product size, product design, and volume. In general, the mold is made up of two components: the two holding plates (top and bottom) and the mold cavity. The larger the platen size, the larger the mold size and the more cavities per mold. However, if the molds are designed correctly (holding plates and cavities), the same cavities can be used in the different size holding plates. This is a major issue in today’s environment where flexibility is essential to be successful. As mentioned earlier, to manufacture lip type rotary shaft seals with a molded primary lip to a consistent highquality standard using the compression molding process, it is necessary to use the “shelf” loading technique for the elastomeric compound preforms.

Bond

460

t

Figure 20 A horizontalreciprocatingscrewinjection molding machine. !

9 ; two “bumps” are applied between 5 and 10 S after the mold initially closes. The term “bump” is used when the mold is opened fractionally, approximately 1 or 2 mm for a fraction of a second, to allow all the gases to escape from the mold.

Figure 18 A typical compression press molding line.

2. Injection Molding Rubber injection molding machines used to manufacture shaft seals fall into three categories: Horizontal reciprocating screw machines (Fig. 20). Vertical screw/ram machines (Fig. 21). Vertical reciprocating screw machines (Fig. 22). In recent years, the trend has been toward an increasing use of the reciprocating screw machines. One of the most commonreasons given for the change isthat the “first in-last out” principle applies for screw/ram

Figure 19 illustrates one type of compression mold suitable for molding lip type shaft seals with a metal outside diameter (O.D.). Figure 19a shows the compound preform resting on the “shelf” and the mold beginning to close. The mold continues to close until it gets to the position shownin Fig. 19b. Finally, the mold opens as shown in Fig. 19c.It can beseen that this molding technique produces a molded lip seal with the flash alreadyremoved.In general, whenusing a vacuum chamber on the press, to ensure that all the air and/or gases are evacuated from the mold after closure, one or

Metal

insert

Metal Insert

I

Rubber Preform

Mold open before molding process (a)

Mold closed

Mold open after molding process

tb)

Figure 19 Compressionmoldingprocess.

(Cl


461

Injection molding machines vary. in size, but those most commonly used for shaft seal manufacture range in size as: Platen size size Shot Load

Figure 21 A screw/raminjectionmolding machine.

machines. The major difference between the two types of machines is as follows. Reciprocating Screw Machine The rubber compound is fed into the machine and the screw loads a known volumeof rubber into the cylinder. When the cylinder is completely filled, the screw then acts as a ram to force the rubber compound into the runner system and the mold. ScrewlRam Machine In general, their are two components, a screw and a cylinderham assembled in a 90” “V” configuration. The rubber is fedinto the screw, which forces the rubber into the cylinder. Whenthe cylinder is full, the ram then forces the rubber into the runner system and the mold.

..... .. . ..

Figure 22 A vertical reciprocating screw injection mold-

ing machine.

250 x 250 mm to 500 x 500 mm 50 to 220 cc 800 to 3000 kN

To ensure consistency, the two critical areas for control specificto injection moldingare shot size control and temperature control on the injection system, i.e., screw portionfor the reciprocating screw machines, and both the screw and ram sections for the screwham machines. As mentioned previously, for temperature and cure control, the use of microprocessor controls has significantly improvedthe performance of injection molding machines, including control of the shot size. An alternative approach for injection molding is the use of a rotary table injection machine.This type of machine uses the multistation concept, each station having a single cavity mold generally using a “point” gating/ injection point as previously mentioned. Consider a rotary table machine with eight stations, station one would be the injection station and station seven would be the unload and load metal insert station. This type of process has been successfully used for different types of lip type rotary shaft seals where more expensive elastomers are used and waste needs to be kept to a minimum. The trend toward automation hasresulted in the increasing use of the vertical injection molding machines as shown in Fig. 22. This type of machine has horizontal platens and, as a result, the mold is in the horizontal plane making it easier to load and unload when using automated procedures, including “pick” and “place” robots. Figure 23illustrates a cross section of halfa fourcavity injection mold used to produce rubber covered lip type shaft seals. Figure 23a shows the mold in the closed position and Fig.23b shows the mold in the open position highlightingthe seal (component),the flash cap, and the compound runner. It can be seen that in this case, the “diaphragm” gatingsystem has been used and the seal being removed after molding has had the flash removed during the molding cycle. Figure 24 illustrates the configuration of the four cavities in the mold as shown by the seals as they come out of the mold, including the flash caps and the runner system. Over the years the trend has been towardan increasing use of the more expensive elastomers. Concern has been expressed about the cost of waste compound when using a multicavity mold and injection molding, primarily due to the amount of elastomeric compound in the runner system. There are three approaches to overcome the problem:

1. The use of smallinjectionmoldingmachines with a single cavity mold. This approach removes the need for a runner system and, whenusing a “dia-

Bond

462

ner system can be used to fill the mold during the next molding cycle. Therefore, the onlywastecompound generated during the molding process is the flash cap, and, for rubber covered seals, a small amount on the outside diameter of the seal that helps the rubber flow. 3. The use of multinozzle injection machines. In general, thistype of machine hasa horizontal reciprocating screw that feeds four injection nozzles each connected to a single cavity mold thus minimizing waste. However, one concern with this approach is the need to have the mold accurately lined up (the top and bottom surfaces of the mold willbe inthe vertical plane) to eliminate compound leakage at the nozzlehold interface.

Figure 23 Four-cavityinjection mold: (a) mold closed, (b) mold open.

phragm” gating system, the only waste is the flash cap, and, for rubber covered seals, a small amount on the outside diameter of the seal that helps the rubber flow. The single cavity mold approach also lends itself more readily to automation using a “pick” and “place” robot. 2. The use of multicavity injection moldstogether with a “cold” runner system. The objective of the “cold” runner system is to prevent the runner system in the mold from curing during the molding process. The result is that the elastomeric compound heldin the run-

Components

Figure 24 Sedflash caplrunnersystemconfiguration from four cavity mold.

3. InjectionlTransferMolding Injectiodtransfer molding is generally used for high-volume manufacture. In principle, the transfer “pot” replaces the runner system in a conventional injection mold and the transfer pot contains a ram that forces the uncured compound into the individual cavities making up the mold. However, like largeplaten injection molding using high-cavitycount molds, there is an excessive compound wastethat is veryundesirable when usingthe higher cost elastomers. The large runner system from injection molding is replaced by a thin sheet of cured compound from the bottom of the transfer pot. One approach to overcome the compound waste issue has been the development of the “cold pot” system, which has the same result as the “cold runner”

R

Figure 25 Schematicdiagram of an injectiodtransfer moldingmachine [9]. (a)Hydraulicseparationunitfor upper mold plate. (b) Hydraulic separationunit for middle mold plate. (c) Shuttle systemwith automatic sprue nipple removal. (d) Brushing unit for cleaning middle mold plate. (e) Hydraulic ejector for automatic ejection.

463

Elastomeric Compound 4 -

Shuttle

____c

(Hot PlatelCold Pot)

I Hot Plate I -i

Components Figure 26 A typical mold/transfer pot layout, which can be used in the injectiordtransfer process and a typical component (seal) layout.

systemon conventional injectionmolding that minimizes the amount of waste compound generated during the molding process. The use of a “cold pot” injection/ transfer molding process has all the advantages of highvolume compression molding without having to make rubber preforms andthen loading the preforms into each cavity. Figure 25 shows a typical injection press that can be used for “cold pot” injectionltransfer molding. It is important to note that the molding cycle is different than the conventionalinjection process. The major difference is that when the rubber compound has been moved from the transfer pot to the individual cavities, the transfer pot is removed andthe “hot plate” moved in to take its place. Figure 26 illustrates a typical moldjtransfer pot layout that can be used inthe injectiodtransfer process. The layout also shows the “hot plate” ready to move into position when the transfer pot is removed and a typical layout for seals produced on the injectiodtransfer process. It can be seen that there is noflash or waste compound. The small amount of waste is, in fact, removed with the “hot plate.” This system lends itself to high-volume automation. SUMMARY There is no one individual issue that ensures product and manufacturingprocess consistency. However, there are a number of elements that are essential if consistency is to be achieved: 1. The component manufacturer and the raw material suppliers should work closely together to ensure that

appropriate and realistic raw material specificationsare agreed upon.The specifications should be developedto enable the elastomeric compounds produced with those ingredients to be consistent and processable in the manufacturing equipmentused. 2. The manual content involved in the compound mixing operation should be minimized. However, economic issues must be considered and the trend, therefore, is toward computer weighing systems (generally, a combination of full automation for critical ingredients and fillers combined with a computer prompted, manual system for the remainder), computer controlled mixers, and computer controlled blending mills. 3. To ensure that consistent, high-qualityparts can be molded on any of the three molding processes discussed: compression,injection, or injectionltransfer,the following issues must be addressed: a. The temperature distribution across the mold (top and bottom) should be minimized. b. Every effort should be made to have a uniform pressure distribution across the mold. c. It is advisable to use presses fitted with either vacuum chambers or applying vacuum directly through the mold. d. Press control is critical and if one is to ensure that the elastomericcompound reaches the required cure state yet minimize the time the component spends in the mold, it is essential to use microprocessor controls. e. The elastomeric compound formulation is critical since product performance is dependent on the molding process used.

Bond

464

The choice of moldingprocess is dependent on three primary issues: the volume of parts to be molded, the elastomeric compound to be used, and the product design. The injection andinjectiodtransfer processes lend themselves more readilyto total automation. For lower volume requirements, conventional smallerinjection molding machinesare appropriate; when the higher cost elastomers are being used, the preferred options appear to be either the use of single cavity molds using small injection machinesor multinozzle injection machinesor the use of a cold runner system with a multicavity mold. Both methods minimize the amount of compound waste. For high-volume requirements, the injection/transfer process using a “cold pot.” For short production runs, the small press approach appears to be the most commonly used. Tominimizethroughput time, componentinventory, and work in process, it is necessary to use the “pull” system for manufacturing. The object being to have the ability to respond quicklyto customer demands throughefficient, fast response manufacturing techniques, not through supplying from finished goods inventory, which is costly. The “pull” system requires that the suppliers to the component fabricator, the component fabricator, and the fabricator’s customers work closely together. This is essential as there is little room for error that causes costly “line down” situations for the fabricators’ customer. The trend is also for each part of the chain, supplier, fabricator, and customer, to take responsibility for the quality and performanceof their own products. This has resulted in a term in general use today, “self-certified supplier.” It can be seen that in today’s environment where quality andconsistency are expected, the only waythis can be achievedin aneconomic way isfor the suppliers, component fabricators, and their customers to work closely together. Each should keep the other informed of

issues of importanceto ensure a constant flow ofproduct and, where appropriate, work together on joint developments. ACKNOWLEDGMENTS The author wishes to thank Chicago Rawhide (Americas) for permission to present the information contained in this chapter. REFERENCES 1. R. Bond, “Compound Ingredient Consistency-Its Role in

2.

3. 4. 5.

6.

7.

8.

9.

Process Capability,” Ninth Annual Meetingof the Polymer ProcessingSociety,Manchester,England, April 5-8 (1993). J. S. Dick, and H. A. Pawlowski, “Applications for the Rubber Process Analyzer.” Rubber Division ofthe American Chemical Society, Nashville, Tennessee, November 3-6 (1992). R. Bond, J . Appl. Polym. Sci.: Appl. Polym. Symp. 53: 219-232 (1994). “F” Series Banbury Mixers. Farrel Corp. Bulletin No. 224-A. R.Bond and Y . Takada, J . Appl.Polym. Sci.: Appl. Polym. Symp. 50: 175-184 (1992). R. Bond, “Shaft Seal Manufacture-Alternate Molding Techniques.” Polymer Processing Society European Regional Meeting,Prague, Czechoslovakia, September 21-24 (1992). A. S. Berens andJ. H. Born, “The Effect of Surface Microstructure on the Performance of Rotary Shaft Lip Type Oil Seals.” Fourth Rubber and Plastics Conference, July 4-7 ( 1974). L. A. Horve, “The Correlation of Rotary Shaft Lip Seal Service Reliability and PumpingAbility to Wear Track Roughness and Microasperity Formation.” S.A.E. Paper 910530 (1991). The NewITM Injection-Transfer Molding Technique. Maschinenfabrik Dieffenbacher J. GmbH & Co., Germany.

30 Advancement in Reactively Processed Polymer Blends Mohini M. Sain and Claude Daneault University of Quebec, Trois-Rivieres, Quebec, Canada

1.

INTRODUCTION

This chapter discusses the processing technologies involved in preparing rubber-rubber and rubber-plastic blend vulcanizates and reviews recent developments in the analysis and performance of these materials. The capital intensive processes involving synthesis of new polymers with modifiedproperties are becoming less interesting. This isbecause polymer blending has become a widely acceptable method of modification of polymer properties. A review of rubber blends was publishedby Ronald [l]. The emphasis was more on thermodynamics and properties of rubber-rubber blends. Coran [2] also summarized the properties of various rubber-plastic blends with special reference to technological compatibility. In recent years, interest has been generated by polymer blends involving polyolefins and PVC in combination with natural and various new elastomers. Simple mixtures of elastomer and plastic were commercialized to prepare either impact resistant plastics or easily processible rubber [3] blends. Understandingthat with technological compatibilization blends of polymers can have properties better than those of the component polymers, more development in this area has been madein recent years. However, technological difficultiesare frequently faced in some type of mutually incompatible polymers. This incompatibility may bedue to either viscosity difference or thermodynamic incompatibility.These problems are generally encountered with selectivityof elastomers and plastics. Additionof a phase compatibilizeror modification of component polymers usually improves product immissibility by reducing the surface energy dif465

ference. In fewinstances, optimization of the processing parameter improves compatibility. A definite route to achieve compatibility in rubber or rubber-plastic blends is covulcanization or cocrosslinking.Inelastomer-plasticblends covulcanization needs to satisfy stringent requirements of cure rate and inhibitionperiod.Whenblend components are cured with the same curingagents, the equality of cure rate is hardly achieved and mostly poor cocuring results either from the solubility differencesof curatives in each polymer or due to their unequal reactivity [4,5]. The controlled addition of curing additives in some cases give covulcanizates. For example, maleic anhydride modified EPDM showed improved covulcanizate properties with natural rubber (NR). A new process to develop interface vulcanization is grafting of selective accelerators onto a polymer chain, which in the subsequent process of vulcanization acts as an effective cure accelerator for the second polymer component in the blend. Beniskaet al. [6] prepared SBRPS blends wherethe polystyrene phase was grafted with the accelerator for curing SBR. Improved hardness, tensile strength, and abrasion resistance were obtained. Blends containing modifiedpolystyrene and cis-l &POlybutadieneshowedsimilar characteristics as SBS triblock copolymers. Another area of recent interest is covulcanization block in copolymers, thermoplastic rubbers, and elasto-plastic blends by developing an interpenetrating network (IPN). A classical example for IPN formation is in polyurethane elastomer blended acrylic copolymers

VI.

We have found that proper choice of curatives and the effective modification of polymer components im-

Saln and Daneault

466

proved covulcanizate properties in polymer blends.The focus of this chapter is on morerecent advances in covulcanization of modified and unmodified rubber-rubber and rubber-plastic blends.The modification process adopted here is to add a controlled amount of special cross-linkingagents before the final moldingprocess, so that the two phases mutually interact to form chemical linkages between them. Such physically mixed blends are expected to resemble block copolymers and would show improvement of some properties of semi-rigid plastics, particularly, impact strength and yield stress, as well as improve elasticity of softer grade thermoplastic rubber vulcanizates. This chapter will deal with the methods of preparing these modified elastomer-plastic and rubber-rubber blends.

II. TECHNOLOGY FOR SOME PROCESS REACTIVE BLENDS PREPARATIONS The processingtechnologies for elastomeric blends, thermoplastic elastomer-based on mechanical mixing, and elastomer-plastic vulcanizates are distinctly different. Depending on the type and nature of blend, size, and their final application, a wide range of processing equipment is now in use both industrially as well as in laboratory scale preparation. A.Elastomer-Elastomer Blends with Dlssimilar Cure Rate The oldest technology involved inthe elastomer blending and vulcanization process is essentially a temperature controlled two roll mill as well as internal mixers followed by an optimum degree of crosslinking in autoclave molds (compression,injection, etc.) in a batch process or in a continuous process such as continuously heated tube or radiated tubes. A few examples of laboratory scale preparation of special purpose elastomeric blends is cited here. Rubber blends withcure ratemismatch isa burning issue for elastomer sandwich products. For example, in a conveyor belt composite structure there is always a combination of two to three special purpose rubbers and, depending onthe rubber composition, the curatives are different. Hence, those composite rubber formulations need special processing and formulation to avoid a gross dissimilarity intheir cure rate. Recent research in this area indicated that the modificationof one or more rubbers with the same cure sites would be a possible solution. Thus, chlorosulfonated polyethylene (CSP) rubber was modified inlaboratory scale with 10 wt% of 93% active meta-phenylene bismaleimide (BMI) and 0.5 wt%of dimethyl-di-(tert-butyl-peroxy) hexane (catalyst). Mixing was camed out in an oil heated Banburytype mixer at 150-160°C. The addition of a catalyst was very critical. After 2 min high-shear dispersive melt mix-

ing of CSP and BMI, the catalyst was added and the mixing reaction was continued for another 3 min. The system temperature was controlled carefully to avoid any possible increase in temperature. The composition was then dumped andchopped into pieces to use further as a reactive cure site. This BMI-treated CSP composition is designated as “modified CSP” (MCSP) in the text. All dissimilar elastomer compositions, whether unmodified or modified, were prepared in a Banbury mixer at 120-140°C. The recipes of unmodified masterbatches are given in Table 1. Four recommended and generally used curingsystems were used for three selected elastomer blends. When ionomeric MgO cross-linking systems were widely used for separate crosslinking of CSP and CM rubbers, this same system was used to understand the features of cure-rate incompatibility in the CSP-CM blend. A typical curingsystem recommended for acrylic rubber only wasalso taken as an example to explain the need for high selectivity of common curing agents in elastomer blends. The compositions for the curing systems are given in Table 2. B. Preparation of Elastomer-Plastic Blend by Melt Mixing Without Vulcanization Elastomer-plasticblendswithout vulcanization were prepared either in a two rollmill or Banbury mixer. Depending on the nature of plastic and rubber the mixing temperature was changed. Usually the plastic was fed into the two roll mill or an internal mixer after preheating the mixer to a temperature above the melting temperature of the plastic phase. The plastic phase was then added and the required melt viscosity was attained by applying a mechanical shear. The rubber phase was then added andthe mixture was then melt mixed for an additional 1 to 3 min when other rubber additives, such as filler, activator, and lubricants or softeners, were added. Mixing was then carried out with controlled shear rate

Table 1 ElastomerMasterbatchCompositions

Composition Rubber PEG 300 Stearic acid HAF black FEF black Chlorinated paraffh Ketone-amine condensate

CM-MB

CSP-MB ACM-MB 100

0 10

100 0 1.o 50 0 0

0

0

3.0 2.0 50

100 0 0 0 55 0

0.25

MB-master batch; ACM-acrylic rubber; CM-chlorinated rubber; PEG-polyethylene glycol; MV-ML(4 + l), 100°C for CSP, ACM, and CM masterbatches are 34, 30, and 32, respectively.

Reactively

467

Table 2 RecommendedCuringSystems

Composition 2.0

14/40

MgO DPTS TDD NC B-l8 DOTG Peroxide TAC

a 7.0

-

-

b

C

0

7.0

0.8 0.5

-

-

ac

-

2.5 3.5

-

-

d

ab

-

3.5 1 .o

-

-

0.4 0.25

-

5.0 1.0

-

bc

ad 7.0 1 .o

1.25 1.75

-

7.0

3.5

-

1.25

1.75 0.4 0.25

2.5 0.5

-

-

1.0

-

DF'TS-di-pentamethylene thiuram tetrasulfide; TDD-thiodiazole derivative; NC-fatty acid amide amine; B-18-special curative; DOTG-diortho-tolyl guanidine; peroxide 14/40-dicumyl peroxide; TAC-triallyl cyanurate.

and temperature, and the mixture wasthen dumped into a granulator to form granulates for further use. The difliculty of the system is the temperature sensitivity of the compound. As soonas the mixing is over it is essential that the molten massbe released fromthe mixer without much delayto avoid hardeningof the plastic phase. The granulated product is then either compression molded or injectionmolded to obtain a test specimen. The compression mold as well as injection mold should be provided with heating and cooling arrangements and the products must be cooledat least 40°C below the melting point of plastic phase to avoid stress cooling. C. Preparation of Elastomer-Plastic Blends by Static Vulcanization

NBR-PVC blend was one of the first invented statically cured elastomer-plastic systems to find industrial application [g]. The concept of vulcanized elastomer-plastic blend was introducedeven earlier by Fischer [9]. Insuch blends, the rubber phase is crosslinked by conventional curing systems. However, crosslinkingis carried out only after a medium- to high-speed shear mixing of the molten plasticphase with the heat soften rubber phase. We have madeseveral statically cured elastomer-plastic blendsincludingNBR-PVC,EVA-PE,PRP-EVA, NR-PP, NR-LDPE compositions. The blend preparation process is very similar to that used for unvulcanized blends. However, in this case respective curing agents are also added in the internal mixer 1 min prior to the unloading process. However, the mixing temperature is always keptat least 20-30°C lower than the curing temperature. Generally, a temperature of about 100-140°C is used for melt mixing, depending onthe curative. For conventional sulfur and peroxide curable blends, the mixing temperature isusually restricted to 120"C, whereas for high-temperature curable blends the mixing temperature can go as high as 140°C. For both single phase crosslinking and cocrosslinking the curatives were added just before the end of the mixing cycle to avoid

premature vulcanization or covulcanization. After the mixing process the compositions are cooled and then granulated. They are finally compression molded for a giventime at their optimum cure temperature determined bya Monsanto rheometer. Table 3 gives the stock mixing and molding temperatures. D.

Elastomer-PlasticBlendsbyDynamic Vulcanization

Gessler [101 prepared semi-rigid rubber-plastic compositions containinga small amount of vulcanized rubber. Later, Fischer [9] and, more recently, Coran and Patel [l11 used the same techniques to develop partially vulcanized blends, which are thermoplastic in nature with different end use properties. Coran and Patel [l21 used various combinationsof plastics andrubber with a large number of curing systems, and they have concludedthat if enough plastic phase is present in the molten state, then the compositions are processible as thermoplastics. A similar methodas used by Gessler [lo] and Coran and Patel [l21 was usedin our study. The elastomer-plastic compositions were different from those described by others in that the thermoplastic part contained a highly nonpolar polyolefin phase anda highly polar vinyl polymer. The elastomers were either naturally occurring hy-

Table 3 MixingandCuring Temperatures

Curing Blend composition NBR-PVC EVA-PRP NR-LDPE NR-PP NR-PRP

Mixing temperature ("C)

110 180-190 120 170 180

temperature ("C) 160-180

200

150-160 180 170-190

.

468

Sain and Daneault

drocarbon polymers or synthetic elastomers containing polar pendantgroups. The process used wasas follows. All preparations were carried out either in a Brabender internal mixer only or by using a combination of a Babury mixer and a twin-screw extruder. For blends prepared in the Brabender mixer, equipment wasfirst preheated to the melting temperature of the plastic phase. The plastic phase was then added and mixing was continued at about 40 to 60 rpm for several minutes until the plastic phase melted completely.Rubber was then introduced slowly by applying rampressure and mixing was continued for 2 min to disperse the plastic or the elastic phase, whichever wasless in quantity. Other compounding ingredients were added slowly including compatibilizers, activator, fillers, and plasticizers and mixing was continued for an additional2 min. At thispoint vulcanizing agents were added and the mixing speed was increased from 60 to about 90 rpm and mixing was continued untilthe torque began to decrease. The batches were then removed fromthe mixer, pressed, and sheeted out to approximately 3-mm slabs. Final molding was done by using the following cycle at a given temperature determined by the type of composition: preheating for 10 min, pressing between the platens for 15 min, and cooling for 15 min at 60°C. In another case where the twin-screw extruder was used, the rubber and plastic were melt mixed with all ingredients in a similar manner as described in blend compositions for static vulcanizations. The product was then dumped, cooled, andgranulated. The premixed granules werethen fed into a twin-screw extruder where a very narrow temperature profile was maintained with a relative high compression (2: l), and the screw speed was adjusted depending onthe final torque and the flow behavior of the extruded stock. The stock was cured by shear force and temperature enforced bythe twin-screw extruder. The dynamically crosslinked blend was taken out in the form of a strip or solid rod to determine the

physical and mechanical properties and the flow characteristics. 111.

PROPERTIES OF REACTIVELY PROCESSED BLENDS

Unlikeincompatible heterogeneous blends of elastomer-elastomer, elastomer-plastic, and plastic-plastic, the reactively processed heterogeneous blends are expected to develop a variable extent of chemical interaction. For this reason the material properties, interfacial properties, and phase morphology of reactively processed blends would differ significantly from heterogeneous mixtures. A.

Properties of Elastomeric Blends with Dissimilar Cure Rates

Table 4 demonstrates the cure rate differences among different elastomers with their respective recommended curing agents. Virtually, there is no match in the optimum cure time betweenCSP and CM or ACM. It is then obvious that a combination of respective curing systems for elastomersin a two phase blend willresult in overcuring of CSP and/or undercuring ofCM or ACM. Two different cure systems were tested on CM rubber. However, none of them appeared to be a closer cure match to unmodified CSP. Similarly, the mechanical properties of CSP-CM blends are extremely poor compared to that of CSP. Interestingly, it was foundthat the cure system “ab” was an excellent combination to achieve a good cure of the CSP-ACM blend as it exhibiteda synergism in torque value. Some important mechanical and heat aging properties of those blends are given in Table 4. Even the CSP-ACM combination gave a high torque at optimum cure, but it failed to attain a good strength property. Apparently, the nature of chemical bonding

Table 4 Cure Characteristics and Properties of Unmodified Blends

Stock

CSP-MB ACM-MB 50 50 CM-MB m Cure emp. 155 Cure 155 me 30 Cure 9 22Max. tor. 30 12.7 Tensile 15.8 st. 11.8 740 Elongation 684 set Tension Hardness 67 Aged, 150°C 72 h, %reten 4Elong. 12

2 7

l 6

5

3

4

100 100 165 20 Nm 41 MPa 12.4 375 % %

Sh 61 A

Tensile 50

50

50

ab

ad

100

Phr

100

a 155

b

C

11

30 28 16 24

26 29

37 23.7 244

1 .o

81 153 52 4

4.5 74 90 110

2.0 78 163 69

d

2.0 72 80

bc

1.5

1 .o 10

4.0 59 81

e

Reactively Processed Polymer Blends

469

during crosslinking of CSP-ACM blends is not identical to those formed during theirseparate vulcanization process. This is because the strength property improved dramatically after aging at 150°C for 72 h. It is possible that some slow intermolecular reaction continues between the two elastomer phases during heat aging of the blend, which makesthe blend more resistantto thermooxidative degradation. No proof of such a reaction, if any, was documented. In contrary, the cure rate mismatch was very significantfor the CSP-CM blend system. An incompatible CSP-CM system shows additive values of mechanical, and aging properties. No synergism was found in mechanical properties. The modification of CSP with inbuilt additional cure functionality showed abetter curecompatibility withthe CM system. Table 5 furnishes the cure characteristics and some useful properties ofcure compatible blendsafter functional modification ofCSP (MCSP). It is evidentthat theoptimum cure time of MCSP is closer to that of the CM rubber. Moreover, a50:50 blend of MCSPand CM rubber has a cure time ofabout 20 min, which is reasonably close to that obtained for each of the rubber components. The cure modified blends showed improved oil and heat resistant properties compared to a blend with high cure rate mismatch. Another advantage of BM1 MCSP was its improved scorch safety. It is predicted that the complex degradation reaction of sulphonyl chloride groups generates free radicals to react with BMI. This MCSP rubber apparently bound with pendant reactive functionality and probably retarded the cure rate. It is further believed that the better cure compatibility gives rise to a better distribution of crosslink density in each rubber phase resulting in a better heat and oil resistance.

1 3.5

5[

2.5 2

g

z

W-

3

g 0

1.5 t, a 1 :

0.5

0

a

20

1m-

40 60 LDPE Loading, wt%

Figure 1 LDPE loading on inhibition period torque; 1,2peroxide; 3,4 = BMI. Source: Ref. 14.

are shown in Fig.1.Evidently, increasing the polyethylene load decreases the rheometric torque monotonously for the BM1 curing system, and for the peroxide cure system, it passes through a minimum. Moreover, the measured torque was always higher with peroxide.It is known that in the presence of peroxide, the LDPE phase also gets crosslinked in addition to the NR phase. It is also reported that BM1 does not cure PE [13]. Thus, the lowering of torque with BM1 is explained by the plasticization effect of PE and the changes oftorque with LDPE in the peroxide cured system is a complex dependence of plasticization effect ofPE and degree of crossB. Properties of Conventionally Vulcanized linking of both the phases. From Table 6 , it is evident Elastomer-Plastic Blends that the optimum crosslinking depends ina complicated I . NaturalRubber-PolyolefinBlends manner on the cross-linkingrate. Thefirst order vulcaniThe cure characteristics of NR-LDPE blends containing zation rate constant ( k ) is higher for pure rubber than that of pure LDPE in a peroxide cured system. Furthervarying amounts of LDPE for two different cure systems Table 5 Cure Behavior and Properties of Modified Blends 1 9

Stock CSP-MB MCSP-MB CM-MB Cure Op. 20 32 Max. 22 aged Heat at 125"C, 168 h ASTM Oil 3 aged at 100°C and 72 h

4

100

Phr Phr Phr

50 100

23

20

11 37 % retention tensile 123st. 121 109 106 116 % retention elongation 77 72 % retention tensile st. 78 67 % retention elongation 72 70

16

7

10 100

50

50 50

30 89 62

70 74 59 62

69 71

68

Sain and Daneault

470 Table 6 Vulcanization Rate Constant (k) and Optimum Cure

Time (790) k

NR-LDPE (DCP) Composition

TW

min

(DCP)

100:o 90:10 65 :35 0.366 50: 50 0.322 0.201 15.5 0.282 10.4 35 :65 10:90 0: loo 0.277

8.5 7.3 8.1 9.1

0.309 0.2017 0.383

11.0

0.272

(BM11

k (BMI) min"

17 16 16 15.5

0.181 0.189

-

-

7%

min"

0.165

10.7

rnin

DCP = dicumylperoxide;BM1 = m-pheneylene bis-maleimide. Source: Ref. 14.

more, a vulcanization rate synergism has been achieved with low PE content in the blend. A probable explanation of such a synergistic effect is the occurrence of an additional cross-linking reaction besides the crosslinking of the two individual phases. Depending on the blend composition, this additionalreaction seems to be possible, assuming a large interpenetrating interfacial area formed during peroxide curing of the blend where cocrosslinking can take place at the interface through peroxidelinkages. Thus, in contrast to elastomer-elastomer blendswhere a microphasecovulcanization is best characterized by the activation energies and the induction periods rather than by the vulcanization rate [14], the vulcanization rate probably plays a dominant role over the induction period. Evidently, the mechanical and processing properties of the conventionally vulcanized NR-LDPE blend would be guided bythe nature of the curing agent. Table 7 gives some important properties of the NR-LDPE blend containing65% elastomer phase for different curing systems. The reactivity of the sulfur system was very much restricted to the rubber phase only. BM1 reacted preferably with the rubber phase as a curing agent, but

at the same time it modified the LDPE by in situ free radical reaction withthe formation of some weakinteraction between cured rubber and the LDPE phase [13]. Such interaction, if any, would leadto better mechanical strength and high mixing torque. This is evident from Table 7. It is evident that neither sulfur nor the BM1 system shows a tensile synergism (T.S. of NR = 12.6 MPa and of LDPE = 15.8 MPa). Peroxide vulcanized NR is seldom used commercially due to its poor mechanical properties [15]. Our earlier report [l51 suggested that by optimizing the NR-LDPE ratio, cure additives, and stabilizers, a blend with mechanical properties comparable to carbon black reinforced NR vulcanizate can be achieved. It seems that the tensile synergism isa direct effect of crosslinking of the plastic phase. However, interfacialadhesion strength of the NR-LDPEblend cured with the peroxide-HQ system was found to be much higherthan the correspondingsulfur cured system [15], suggesting a cocrosslinking of NR with LDPE.The disadvantage of such blends are their poor thermal stability. In fact, a peroxide cured blend could be easily degraded under an accelerated aging condition. Therefore, a suitable choice of antioxidant was found to be

Table 7 Effect of Various Curing Systems on NR-LDPE (65:35) Blend

Curing Stearic System

SNBTSIZnOl Acid

310.8 Composition Property 33.6Torque (Nm) Mod(3) (MPa) MPa) 13.9 Tensile 13.6 Hardness (A) 22 Tenset (%) Elong. 510 (%)

2/1.31511 22.9 2.7 7.9 16.1 48 24 510475

BMINBTS 5.010.6

DCP (40%)

DCP/HQ DCPlSulfur DCPIEGDM

3.510.35 5.0

8.7 13.4

11.2

65

63

400 550

415

EGDM-ethylene glycol dimethacrylate; HQ-hydroquinone; Mod(3)-300% Modulus.

2.015.0 8.9

9.7

14.7


471

l

11

18 I

essential to improve the aging behavior of peroxide cured NR-LDPE blends. The best recommended antioxidant system isa suitable combination of 1,1,3-tris(2methyl-4hydroxy-5-tert-butyl phenyl)butane, dilauryl thio-dipropionatewithDCP. The aging resistance of such a composition wasgreater than 60% at 100°C after 72 h. The other advantage of blending NR with PE is the improvement of electrical properties of the relatively poor electrical insulant, NR. Figure 2 shows the effect of the LDPE content on volume resistivity ( p v ) and loss tangent (tans) values of the blends. The increase in the LDPE loading causes (1) a sharp increase in volume resistivity, and (2) a decrease in the loss factor. Blend permitivity remained practically unchanged, whichis about 2.4for any givenvulcanizate containing minimum 10% PE and maximum 9 0%PE. Unlike LDPE or HDPE, other polyolefins are difficult to crosslink by the peroxide curingsystem because polyolefins containing propylene as a monomer unit in the main chain are very prone to free radical decompositionwith peroxide. Therefore, in a elastomer-plastic system containingPP or PRP (propylene-randomethylene propylene-propylene) copolymer as plastic phase peroxide system is not a successful reaction initiator. Technological compatibility was needed to process PP or PRP with NR inorder to achieve better strength and processibility. Coran andPatel [l61 foundthat phenolic resin could be usefulto modify PP in order to improve the technological compatibility of PP with other polar elastomers, such as NBR. A similar processing technology as described by Coran and Patel [l61 was used to make PP or PRP as a curative for NR. It is well known that phenolic resin can vulcanize NR under certain experimental conditions. Therefore, a PP or PRP matrix containing pendant phenolic functionality would serve as a curative for NR-PP or NR-PRP blends for a result-

16

- 0.8 - 0.6 Lo

-

-

5

0.4

- 0.2 2

"

0

"

" 20

t

' " 40

'

PE LOADING, W

Figure 2 Electricalproperties Ref. 13.

0 100

80

60 e

of NR-LDPE. Source:

ing cocuring effect during the processing or molding operation. In our study [171, both PP and PRPwere treated with 2-4 parts of phenolic curatives in the presence of SnCl2.2H20 and MgO to obtain phenolic modified PP or PRP. The modified polyolefin was then melt mixed with NR under an optimumshear condition and temperature to develop chemical bonding between NR and polyolefinviaquinone methide formation [IS]. About l min before dumping, the batch curatives were added to the mixer to disperse the curatives in a partly reacted blend. The total curing of the rubber phase in a compression mold then improvedthe properties of the compatibilized blend over unmodified PP or PRP blended and cured stocks. Table S gives the properties of the blends after

Table 8 Properties of NR-PP or PRP Blends (Modified with Phenolic Resin) at 50:50 Ratio

Preparation

NR-PRP

Cure system ("C) temperature 180 Cure 180 180 180 time 20 Cure 20 (min) 20 Processing Properties Modulus, 100% (MPa) 9.0 3 29.3(MPa) strength 24.7 Tensile 26.9 23.3 22.5 23.6 Elongation 480at Br.475 (%) 475 88 Hardness 88(Sh A) 86 3 set 26 Tension 28 (%) 27 170 176 (MPa) True 142break stress 135 at 135 129 136 density Cross. 7.5 (l@) 5.8 11.8 Reprocessibility

180

R.M.

= roll mill; B.M. = Source: Ref. 19.

NR-PP

E

E

F

F

E

E

F

F

R.M.

B.M.

R.M.

B.M.

R.M.

B.M.

R.M.

B.M.

180 20

11.9 20.8 400 87 28 104 9.0

Brabendermixer; E

11.13.6 1

9.0

9.911.7 400

475

37.5 28

13.3

500 37.5

No =

DCP-HQ/4.0:0.5%; F

=

PF resin-SnClz/S.O:O.5%.

500

No

Sain and Daneault

472

Temperature ‘C

Figure 3 tan8 versus temperature plot forNR-PP blends (50:50); curve 1,purePP; curve2, dynamically cured blend; curve 3, conventionally cured blend. Source: Ref. 19.

processing in two different machines.The vulcanizates have excellent mechanical properties. The mixing techniques, however, have no significant influence on the blend properties. The blends are not repressible. As mentioned before such blends are technologically compatibilized by reactive processing. The improved phase interaction of the conventionally vulcanized system is apparent from spectra shown in Fig. 3. A shift in the glass transition temperature of NRand the distinct shape of the original sharp PP glass transition peak confirms better interaction for the resin containing the conventionally cured blend.

2. NBR-PVC Blends Our study was also extended to obtain the effect of a selectivecuring system onNBR-PVC compositions.

NBR-PVC is a very useful blendfor the electrical cable sheathingapplication that demands highflexibility, strength, oil resistance, and low-temperatureresistance. Table 9 summarizes someresults of the effect of curing systems during blend processing for a 60 :40 NBR-PVC blend. A marked change is visible on the mechanical properties depending on the curing systems. While the systemcontainingPETA-DHBP develops the best strength properties beforeaging,whichis associated with the high gel content of the vulcanizate, the mixture containingPNDAshowedgood strength only after aging. It is expected that this change in the mechanical properties in the presence of different curingsystems is due to a change in the vulcanizate structure during the molding process. The peroxide covulcanization of NBR-PVC showedpoor thermo-oxidative properties when sulfur curingof the same composition resulted in a single phase cure only. Based on our previous experience, we have also studied the effect of selective cross-linking systems on NBR-PVCblendscontaining dominant PVC phase, 60% by weight. Some experimental results are in Table 10. Evidently, at a higher concentration of BM1 the gel content was somewhat higher than for the gel content prorated for NBR content in the blend. It seems a partial crosslinking ofPVC or NBR-PVC interface has occurred. The results are in conformity withthe mechanical properties of the BM1 cured blend. In mostof these cases the measured values of mechanical strength has gone through a maximum, indicating a change in the reaction mechanism duringthe molding process with curatives. In order to support our prediction that the change in mechanical properties with different curing systems is due to a change in the vulcanizate structure, results werecomparedfromdynamic-mechanicalmeasurements as shown in Figs. 4 and 5.

Table 9 Vulcanization Characteristics and Propertiesof NBR-PVC Blends

Curing system Cure characteristics Temperature (“C) Time (min) MV,ML(l + 4) Properties Modulus, 100% Modulus, 300% Tensile strength (MPa) Elongation (%) Hardness (ShA) Gel content

Source:

Ref. 20.

Sulfur

SulfurPNDA

BMIMBTS

PETA/DHBP

160 5 36

160 5 35

180 15

180 20

1.8(3.7) 3.7(6.5) 7.1(13.3) 740(700) 54(55) 56(69)

1.7(2.8) 3.2(6.6) 8.6(4.5) 780(638) 56(58) 67(84)

3.2(4.3) 5.3(6.8) 6.9(7.7) 475(375) W68) 35.7(42)

6.6(8.7) 17.1 17.7(16.6) 330(215) 80(78) 93(95)

-

-

Reactively

473

Table 10 Effect of Short-Time Aging on Mechano-Chemical Properties of PVC-NBR Blend

SITMTIMOZPNDAI (wt%) system Cure

BMI-MBTS/S.O: 1.7

PETA-TEA 1 0 ~ 0 . 2

8.3(14.5) 15.4(18.6) 16.3( 16.3) 362(282) 75.6(65) 71(69) 54( 10.1)

8.2(13.1)

Modulus 100% (MPa) Modulus 300% (MPa) Ten. strength, (MPa) Elong. at break, (%) True stress at break (MPa) Hardness (ShA) Gel Content (%)

0.1 :0.33:0.44: 1 . 1

7.1(14.3) 13.3(20.3) M.O(l9.8) 405(385) 75.1(95.4) 65(73) 24( 34)

-

11.5(12.8) 255( 173) 41(34.5) 73(73) 52(71)

TEA = triethylene amine. Source: Ref. 20.

The values of gain modulus, loss modulus, and tan8 show only one relaxation peak within the studied temperature range. A comparison of the maximum absorption peaksin Fig.4 indicates glass transition temperature changes withthe curing system, and the relative change is maximum for the PETA-DHBP system. This high T, of the PETA-DHBP system is associated with poor fatigue resistance of the blend vulcanizate [21]. While the PNDA-S system shows maximumfatigue resistance with lowest T, of the blend.

C. Properties of Unvulcanized Elastomer-Plastic Blends

1. Unmodified PRP-EVA Blend In the preceding sections, our discussion has been limited to softer grade elastomer-plastic vulcanizates. Commercial interest, however, also centers on another major familyof polymer blends, semi-rigid impactresistant polyolefins. Thus, we report some of our findings on PRP triblock copolymer and EVA rubber blends without

rubber phase vulcanization at first and then modified interfacial reaction. As withNR-PP or NR-PRP blends, EVA-PRP blendis also incompatible. Therefore, technological compatibilitywas a must to obtain useful properties from this blend. In this case a reactive surfactant was introduced in the EVA-PRP blend duringthe processing operation. This reactive surfactant decreases the surface energy betweenthe EVA and PRPinterface by forming a chemical bond with one of the polymers and physically compatibilizing with the other phase. The interfacial surfactant was prepared by several methods. Inone method used previously by Muras et al. [22], powdered PP was reacted with maleic anhydride (MAH) in solid phase in the presence of a peroxide catalyst below the melting temperature of PP. The grafted polymerwas recovered by extraction and used as an interfacial surfactant for EVA-PRP blend preparation. The interfacial surfactant isdesignated here as MAH-PP.Depending on the grafted MAH content, they are classified as MAH-PP-

- 10‘0 N/m2

- 109

- 108 - 107

-50

-40 -30

-20

-10 0 10 Temperature

20

30 ‘C 40

Figure 4 tan8 versus temperature forNBR-PVC blends; curveI-BMIMBTS;curve2-SPNDA;curve 3 = DHBP/ PETA. Source: Ref. 21.

1

1

1

1

1

1

1

1

-50 -40 -30 -20 -10 0 IO Temperature

1

1

1

20 3O0C4O

1

Figure 5 Gain ( E ‘ ) and Loss (E”) modulus (see Fig. 4). Source:

Ref. 21.

ts

Sain

474

Table 11 Effect of the Mixing Temperature and Processing Method on Pm-EVA Blends ~

~~

Value of the property for the followingEVA content and mix temperature 8 wt%

Property

195°C

30 wt%

170°C

12.0(11.3) Tensile st., 8.3 MPa 18.9(19.5) Elong. at 75(60) % Hardness, 64(65) ShA (7.0) MFI, 5.2 g/10 min 5.7(6.2)

break,

Brittle point, “C Yield stress, MPa18.7(

17.7 16916 4863 5.8 - 50 19.0

47 wt%

170°C

195°C

170°C

7.4

196(82) 53(57)

66

56

- 56( - 50) -)

- 23

8.77.7

-28(-31) 12.3

- 21

9.4

Source: Ref. 23.

I (0.98%), MAH-PP-II(l.41), and MAH-PP-III(2.95%) [23]. Table 11 gives the result of unmodified EVA-PRP blends at variable process conditions. The EVA polymer used wasa rubber grade with 40% vinyl acetate content. The results in Table 11 for unmodifiedanduncured blends showed that blend composition has a maximum influence onthe properties and inthat case the optimum compositionwas 8% EVA.Generally the blends are weak in nature but are flowable. Highertemperature of mixing has a marginal effect on blendproperties and the mixing method hardly influencesthe properties. Apparently the blends are grossly incompatible.

2. Reactivity of Maleic Anhydride Grafted PP on PRP-EVA Blend To introduce some interfacial physico-chemical linkage between EVA and PRP, blends were made by adding different quantities of MAH-PP. Some results are demonstrated in Table 12. The physico-mechanical properties of the PRP-EVAcompositionsmodifiedwith MAH-PPshowed that properties are influenced by MAH-PP concentration. Compositions with better impact strength and improved brittleness can be prepared by varying the modifier concentration. Tensile strength and elongation are not significantly influenced by the addition of a modifier. An increase in the modifier con-

centration is markedby an initial increase in low-temperature impact resistance and lowering of the brittle point of the studied composition. However, bothimpact strength and brittleness temperature deteriorated when the modifier level was very high. Toelucidate the cause of this complexdependence of low-iemperature brittleness and impactstrength on the modified concentration, we have studied the dynamic mechanical properties as well as the flow behavior of unmodified and modified compositions. Figure 6 gives tan8 versus temperature correlation for various blends. The lowering of the glass transition temperature from -45 to -55°C of the PRP polymer in the PRP-EVA composition only suggests a possible phase interaction between the two phases. However, more lightwill be thrown on this subject during a morphological study discussed later in this text. It is also interesting to note that the flow behavior of the blend first marginally deteriorated and then, with further increase in the MAH-PP concentration, it was further improved. It also indicates that at higher concentrations of MAH-PP it might function as a flow promoter.

D. Properties of Dynamically Vulcanized Blends

The method of “dynamic vulcanization” was appliedto mixtures of NR withLDPE, PP homopolymer, and PRP

Table 12 Effect of Concentration of MAH-PP-I11 on PRP-EVA 70:30 Blend

Elongation MAH-PP-I11 Impact Tensile MAH, (wt%) 0.5 1.o 1.S 2.0 3.0 Source: Ref. 23.

st. Content (%) 0.015 0.029 0.044 0.050 0.088

(MPa)

( W

31.0 15.0 51.0 15.7 36.5 16.1 15.8 28.1 15.5

81 79 98 95 88

Hardness Brittle

(g110 min)

(kJ/m2) MFI (ShA) (“C)point 7.7 7.3 7.5 30.1 8.5 8.8

- 36 - 48 - 46 - 46 - 45

64 63 62 63 60

475

Reactively Processed Polymer Blends Table 14 Properties of Dynamically Vulcanized 0.7 -

U.6

-

a5

-

l

-MAY-DOTI

C

,o Q & -

-

0.2

-

Preparation Curing system Cure Temp. ("C) Cure time (min) Properties Modulus, 100% (MPa) Tensile st. (MPa) Break elong. (%) Hardness (Sh A) Tension set (%) True stress at break

Lo

Q3

NR-LDPE (65 :35) Blends

BMIMBTS

SulfurlMBTS

180 5

180 5

4.1 9.5 320 57 12.5 40

0.5 4.0 600 32 45 28

(MW

Permittivity Vol. resistivity (0hm.rn.10"~)

Figure 6 Variation of tan8 with temperature for(1) EVA, (2) PRP-EVA/92:8 blend, and (3 and 4) EVA-PRP blend

92:8 modified with MAH-PP-111. Source: Ref. 23.

triblock copolymer. The main advantage of a dynamically cured blend is the reprocessibility. Therefore, these blends are preferred in many industrial applications where dimensional accuracy and product complexity leads to a high rejection rate. Industrial cables are one example that needs high precision and quality. In many cases such application does not need very high mechanical properties. Thus, various custom-designed thermoplastic elastomers can be used. A suitable blend of polyolefin and natural rubber is goodfor many applications. They are good insulants and are highly resistant to polar solvents, moisture, etc. After dynamic crosslinking, such blends could develop desirable strength without sacrificing processibility. Mixing characteristics

2.35 6.5

2.4 6.3

during dynamic crosslinking of some NR-LDPE blends are given in Table13. The mixing torque decreases with increase in plastic content. The Mooney viscosityof the dynamically cured blend passes through a maximum. However, the dynamically cured blend showed higher Mooneyviscositycomparedwith a statically cured blend. High viscosity may be due to the formation of IPN in the dynamically cured blends as suggested by several coworkers [7,24]. It is predicted that the IPN formation depends on the blend compositions, and once formed, it strongly resists phase separation. This could be the reason for poor flow behavior of the dynamically vulcanized NR-LDPE blendat anintermediate composition as given in Table 13. Therefore, the dynamically cured blend of technical interest was foundto be a composition with35% LDPE content. Table 14demonstrates someimportant properties of this 65:35 NR-LDPE blend for two different curing systems. Evidently, the mechanical properties of the dynamically cured blend are very much influenced by the nature of the curing agent; i.e., the nature of the chemical bondformed dur-

Table 13 Blend Composition and Dynamic Curing Characteristics for NR-LDPEa

Composition MV, torque Dump Mixingtemp. ("C) NR-LDPE 165

90:10 65 :35 50: 35 :65 10:90

50 44.7 a

9.5

162 158 156 155

26

(Nm) 30.4 29.4 27.4 22.6 19.6

Stocks compounded with 3 phr BM1 and 0.8 phr MBTS.

ML[4

+ 13 121

28 44 35

MFI ( d l 0 min) 8.1 1.7 10.2 21.6

P

Sain and Daneault

476

ing the dynamiccuring process. Sulfurcuringgives bonds that are flexible in nature and lead to softer composition. On the other hand, imide linkage formed during crosslinking withBM1 is rigid innature and mechanical properties are reasonably high. Moreover, the blend is reprocessible. Thus, a dynamically cured NR-LDPE blend with BM1 is of commercial interest in electrical industry. In the past several attempts to apply dynamic vulcanization techniques to NR and PP blends [25,26], for both PP homopolymer andPP copolymer, were less effective owing to the difference in their miscibility characteristics. Compositions with practical interest, however, have been prepared by modifying polyolefins and thendynamicallycuringthem. PP homopolymerand PRP copolymer were pretreated with phenolic resin and SnC12 to improve blend compatibility; hydroquinone, a hydroperoxide radical scavenger, was used to prevent degradation of the PP chains. The physical properties of the blends cured by the dynamic process are given in Table 15. It is evident that reactive mixing of NR-PP and NR-PRP could lead to reprocessible thermoplastic compositionswith excellent mechanical strength. A comparison of two curingsystems suggests that phenolic resin curingby dynamic vulcanization probablyleads to a highly compatible and miscible blend withreasonable crosslinking of the rubbery phase. In addition, both these blends haveexcellent electrical properties. Figure 3 demonstratesthat loss tangent value is very reasonable for medium voltageelectrical insulation. However, tan6 versus the temperature of the dynamically crosslinked blend with peroxide (Fig.3) indicates less interfacial interaction between twophases as there is onlya marginal shift in the glass transition values. The morphological changes with phenolic modified blendsare discussed in the following section.

W.

PHASE MORPHOLOGY OF REACTIVE PROCESSED BLENDS

Dissimilar polymer blends show phase separation and poor mechanicalproperties. A suitable processing technique in combination witha technical compatibilization of blend couldreduce the immiscibility significantly. Reactive mixing leading to graft copolymerization or cocrosslinking improved mechanical strength of various polymer blends discussed here. The differences in the mechanical properties between sulfur cured and peroxide cured NR-LDPE blends are very much reflected in their phase structure. Figure 7 shows SEM microphotographs of fractured and etched surfaces of NR-LDPE (65 :35) compositions. The sulfur crosslinked blend has a higher particlesize of the plastic phase, and the dispersion of those particles is also poor. Again, blendscured with peroxide have very good dispersion of the cured plastic phase and the particle size was much smaller. Further morphological investigation of the interfacial adhesion onthe reactively cured blends wascarried out by taking SEM microphotographsunder externally applied predetermined strains. The SEM micrographsof a series of samples with a series of gradually increasing strains have been analyzed. Figure 8 shows the SEM images witha low levelof strain (50%). It is clear that even with a low-strain level defects are initiated in the sulfur cured system with the formation of large cracks at the boundary layer between the two phases. However, in the peroxide cured system the mechanism of crack initiation is very different. In the latter case the NR-LDPE interface is not the site for crack initiation. In thiscase, stress due to externally applied strains is distributed throughout the matrix by formation of fine crazes. Furthermore, such crazes are developed in the continuous rubber matrix in a direction

Table 15 Properties of Dynamically Cured NR-PP and NR-PRP Blends

Preparation Curing system Cure Temp. ("C) Cure time (min)

Properties

Modulus, 100% (MPa) Tensile st. (MPa) Break elong. (%) Hardness (ShA) Tension set (%) True stress at break (MPa) Flow property

DCP-HQ 180

5

7.6 7.6 100 88 34 15.2

Flowable

PF-SnCI2 180 5

10.2 18.5 280 89 37.5 70.5

Flowable

MPP = phenolic resin (PF) modified PP; MPRP = phenolic resin modified PRP.

DCP-HQ

PF-SnCIz

7.5 8.8 150 88 26 22

9.4 16 300

180 5

Flowable

180 10

90

27.5 64

Flowable


477

Figure 7 Optical micrographs for NR-LDPE (65:35) blends: (a) cured with sulfur; and (b) cured with peroxide. Source:

Ref. 19.

perpendicular to the applied strain. This difference in the mechanism of stress distributionis one of the reasons for differences in the mechanical properties of the system studied. A similar morphological feature was also noticed for NR-PPblends. The phenolicmodified NR-PP blends after BM1 and phenolic curing showed a continuity between the two phases. In contrast, unmodifiedNR-PP blends either conventionally or dynamically cured by sulfur chemistry resulted in a grossly heterogeneous structure leading to interfacial separation during mechanical stress or repeated fatigue tests [19]. Effect of processing temperature as well as relative concentration of two phases in a blend were found to have a significant effect on the blend morphology and interfacial interaction. Results of the investigation of PRP-EVA blends with and without maleic modification give more insightinto the morphological properties. The results of an investigation into the unmodified PRP-EVA blend in Fig. 9 illustrate that the mixing tem-

perature necessary to obtain the best blend property is the temperature at which the EVA rubber is best dispersed in the plastic matrix, i.e., the temperature at which the maximum interfacialarea is developed. From an analysis of the SEM photographs and subsequent evaluation of the maximum interfacialarea per unit volume of the blend, it is evident that for both the polymer compositions studied(8% EVA and30% EVA), as illustrated in Fig.9, the maximum calculated value of interfacial area was obtained at 195°C when optimum blend properties were also achieved. The calculated values of the interfacial area from SEM micrographs in Fig. 10 also confirmed that the blend with 8 wt% EVA developed a maximum specific interfacial area with the narrowest particle size distribution. This observation is in line withthe conclusion drawn by Eliott and Tinker [26]. Maleic modified EVA-PRP blend showed some differences in the blendmorphologyfromunmodified EVA-PRP blend. Additionof MAH-PP significantly in-

478

Sain and Daneault

Figure 8 SEM images of etched surfaces of blends with 50% stretching; (a) sulfur cured(3000 x ); and (b) peroxide cured (3000~).Source: Ref. 27.

n

0.35

(30%EVA)

modified EVAlPRP

CCI

.-

0

'c

Q)

- 0.15 unmodified EVAIPRP, 8% EVA U. I

160

170 180 190 200 210 220 PROCESSING TEMPERATURE, C

Figure 9 Effect of processing temperature on interfacial area.

230


Figure 10 Particle size distributionof dispersed EVA rubber in PRP: (1) 1.5% EVA; (2) 8% EVA; (3) 15% EVA; and (4) 30% EVA. Source: Ref. 19.

Figure 11 SEM images of PRP-EVA

479

creased the interfacial area of the dispersed phase as evident from the value shown inFig. 9, indicating a stronger physical interaction between the polymeric phases. Figure 11 shows SEM images of the etched unfractured (a) and(b) unmodified andMAH-PP modified, respectively, and fractured surfaces (c) and (d) fromunmodified and MAH-PP modified PRPlEVA blend, respectively; elastomer domains with a particle size of about 0.5 to 3.0 pm arevisible in the form of dark holes. It is apparent from the fractured surface of the unmodified composition in Fig. l l c that fractures are brittle in nature with “dropping out” (dark holes) of elastomer particles. On the other hand, the modified composition resists phase separation by partly preventing dropping out of the elastomer phase. Dispersed elastomericmicro-particles preferentially “break” (bright spots) with the development of crazes around the broken rubber domain as evident from Fig.

92:s

blends. Source: Ref. 23.

480

Sain and Daneault

4. M. E. Woods and J. A. Davidson, Rubber Chem. Technol., 49: 112 (1976). 5. V. A. Shersnev, Rubber Chem. Technol., 55: 537 (1982). 6. J. Beniska, E. Staudner and E. Spirk, Makromol. Chem. I 161: 113(1974). 0-C-CH, 7. T. Tanakaand T. Yokoyama, Polymer Sci., Part C: PolyII 0L A O 0 mer Symp. 23: 865 (1968). 8. K. Mori and Y. Nakamura, J . Polym. Sci., Polym. Chem. “-PP EVA Ed. 16: 2055 (1978). 9. W. K. Fischer, U.S. Patent 3,758,643, (Uniroyal Inc.) CH, 0 Sept. 11, 1973. l II 10. A. M. Gessler, (Esso Research& Engineering Co.) U.S. -CH~-C-CH-CH,-C-O-CH-~H,~ Patent 3,037,954, June 5, 1962. I I I 11. A. Y. Coran and R. P. Patel, Rubber Chem. Technol., CH, C-0-C-CH3 54: 892 (1981). II \ II 12. A. Y. Coran and R. P. Patel, Rubber Chem. Technol., 0 0 56: 1045(1983). EVA-MAH-PP J. Beniska, and V. Khunova, 13. M. M. Sain,J.Lacok, Kautch. Gummi. Kunstst. 41: (9) 895 (1988). Scheme 1 Source: Ref. 23. 14. M. M. Sain, J. Lacok,J. Beniska, and J. Bina, Acta Polymerica, 40: (2) 136 (1989). 15. M. M. Sain, J. Beniska and P. J. Rosner, Polym. Mater., 6: 9 (1989). 16. A. Y. Coran and R. P. Patel, Rubber Chem. Technol., 56: 210 (1983). 1Id. Thus, it is expected that in the modified PRP-EVA 17. M. M. Sain, I. Hudec, J. Beniska, and P. Rosner, Rubber blend, probably due to interface modification by reactive Chem. Technol.; 61: 748 (1988). processing, a transesterification between the pendant 18. Moms Morton, (ed.) RubberTechnology, VNR, New York (1973). MAH group inMAH-PP, and acetate groups in EVA the 19. M. M. Sain, I. Hudec, J. Beniska, and P. Rosner, Mat. elastomer, as predicted in reaction Scheme 1, the disSci. Eng., A108: 63 (1989). persed rubber particles become more efficient in craze 20. M. M. Sain, J. Oravec, E. Sain, and J. Beniska, Acta initiation. Polymerica, 43: 51 (1992). 21. M. M. Sain, J. Lacok, J. Oravec, and J. Beniska, Acta Polymerica 40: (10) 649 (1989). 22. J. Muras and 2. Zamorsky, Plasty Kaucuk, (in Czech) 28: ( 3 , -137 (1991). REFERENCES 23. I. Hudec, M. M. Sain, and V. Sunova, J . Appl. Polym. 1. C. M. Ronald, HandbookofElastomers, (A. K. Bhowmik Sci., 49: 425 (1993). and H. Stephens, eds.),Marcel Dekker, Inc., New York, 24. D. S. Cambell, D. J. Eliott, and M. A. Wheelans, NR pp. 183-214,1988. Technol., 9: (2) 21 (1979). 2. A. Y. Coran, Handbook of Elastomers, (A. K. Bhowmik 25. W. K. Fischer, U.S. Patent 3,806,558,1974. and H. Stephens, eds.),Marcel Dekker, Inc., New York, 26. D. J. Eliott and A. J. Tinker, tnt. Rubber Conf., Kuala pp. 249-312,1988. Lumpur, Malaya (1985). 3. D. S. Cambell, D. J. Eliott, and M. A. Wheelans, NR 27. I. Hudec, M. M. Sain, and J. Kozankova, Polymer TestTechnol., 9: 21 (1978). ing, 10: 387 (1991).

31 Graft Copolymerization Onto Natural and Synthetic Polymers lnderjeet Kaur and Bhupendra Nath Misra Himachal Pradesh University, Shimla, India

1.

INTRODUCTION

side chains of different polymers.Graft copolymers are generally prepared by free radical, anionic or cationic A polymer isa large molecule built up bythe repetition addition, or by ring opening polymerization ofa monoof small, simple chemical units. In somecases the repetimer onto the backbone polymer. Thus, the graft copolytion is linear whilein other cases the chains are branched mer isa mixture of macromolecules consistingof a block or interconnected to form three-dimensional networks. of constitutional units of type A (backbone) having The polymer can be formed not only throughaddilinear blocks of units of type B (grafts) attached to the backtion, but also through condensation of similar units as bone (Fig. 1). well. In addition to the genuine copolymer macromoleIf a polymer consists of two or more different monocules, the rough copolymer product usually contains mers, then the resulting polymer is called a “copolygrafted backbone and, according to the conditions of mer.” The copolymer isfurther classified into five types preparation, the attendant homopolymer. The part of the depending onthe sequence of monomers inthe polymer graft copolymerthat consists of macromolecules bearing backbone. grafts attached to the backbone will be referred to as 1. Random copolymer-Distribution of the mono“true copolymer.” The synthesis of a graft copolymer mer units does not follow any definitesequence. These requires the formation of a reactive center on a polymer are produced in bulk, aqueous, suspension, or emulsion molecule in the presence of a polymerizable monomer. using free radical initiators of the peroxide type or redox systems. 2. Regularcopolymer-These copolymers have an ordered sequence in the distribution of monomers. These are produced by controlled feedingof the monomers during copolymerization. 3. Alternatingcopolymer-In these copolymers, the monomers are arranged in an alternate sequence. Alternating copolymers are produced by special copolymerization processes. 4. Block copolymer-These copolymers are built I of chemically dissimilar terminallyconnected segments. Ungnfted backbone True graft copolymer Homopolymer Block copolymers are generally prepared by sequential ” Gratt Copdymer Attendant (Graft) anionic additionor ring openingor step growth polymerization. 5. Graftcopolymer-Agraft copolymer consists Figure 1 Components of roughcopolymerproductsin of a polymeric backbone with covalently linkedlateral graft copolymerization. 481

Kaur and Misra

482

Most methodsof synthesizing graftcopolymers involve the use of radical polymerization, although ionic graft copolymerizations are also gaining importance. Graft copolymerization can be carried out either in homogeneous or heterogeneous systems depending onwhether the polymerbeinggraftedissoluble or insoluble in the monomer and inthe solvent being used. Usually grafting can occur by either the “grafting from” or “grafting onto” method as illustrated by the following: Grafting from:An active site generated along a polymer backbonestarts to propagate monomer and thus produces branches.

II. IMPORTANCE OF GRAFTING

Grafting providesa convenient means for modifying the properties of numerous polymers. It is often required that a polymer possess a number of properties. Such diverse properties may not be easily achieved by the synthesis of homopolymers alone but can be achieved through the formation of copolymers or even terpolymers. The formation of graft copolymer with sufficiently long polymericsequences of diverse chemical composition opens the way to afford speciality polymeric materials. 111.

Grafting onto: A growing polymer chain P2 attacks another polymer P,. “PI

+

”PI

-

+ nM (MYn+1(-P2) - P P I + *-P?M’

T

P

(2) I

P2

During formationof the graft copolymer the following assumptions are made: 1. The attachment of grafts to backbones is random [l-31. Each segmentof the backbone has the same probability to be grafted. Segmentis defined as that part of the backbone to which just one graft can be attached. The probability of attachment of another graft to the backbone is not affected by the grafts already attached to the backbone. 2. Copolymer macromolecules are composed of a single backbone having simple grafts attached to it,i.e., the macromolecules are of the comb-like type. No further grafting of grafted chains is contemplated (4).

METHODS OF GRAFTING

The preparation of a graft (-g-) copolymer in most cases starts from a homopolymer A, which is appropriately activated, and then the growth of the graft of monomer B is initiated. The model must fulfill the following requirements. 1. The copolymer is of poly(A-g-B) type.

2. The probability that a graft will grow on a given

unit of the backbone does not depend on the number or size of the other grafts in the molecule or on the size of the backbone to which it is attached.

The efficiency of grafting isdependent upon several competing reactions:

1. Competition between the various species present in the reaction mixture such as monomer, solvent, and backbone for the growing polymer radical, which means that there is competition between chain growth and variouschain transferring steps. 2. Competition for the initiator radical between monomer and backbone. 3. Competitionamong the terminating processes such as disproportionation after the polymer radical has formed. Usually a graft copolymer is a mixture of homopoly4. Competitionamong the terminating processes in the growing graft species. mer, ungrafted backbone polymer, and graft copolymer. If the termination occurs through recombination of The graft copolymer is usually separated from the homogrowing polymeric chains, it should lower the grafting polymer by the solvent extraction method. Often graft efficiency. But even though the efficiencyof this process copolymers prepared fromwater-soluble backbone is low, it has received considerableuse because the reacpolymers (such as starch, gelatin, poly(viny1alcohol) tions are easily carried out. One needs merely to conduct [PVA]) andwater-solublevinylmonomers (such as polymerization of the monomer in the presence of the acrylamide, acrylic acid, etc.) arealso water-soluble. In backbone polymer anda proper free radical initiator. A such cases a rough copolymer is obtained that is not a solvent may or may not be present. The addition of the true graft copolymer but a mixture of ungrafted polymer, good chain transferringagent, such as a mercaptan [8], grafted polymer, and the homopolymer. In such cases CC4or CH31[9], causes reduction in grafting efficiency. homopolymer cannot be removed from the mixture by The other reaction parameters, such as concentration the conventional solvent extraction method. Hence, of the initiator, monomer, temperature, and the time of such products are referred to as composites. Evidence for grafting insuch cases has been obtained from turbidi- reaction, affect the grafting efficiency.The nature of the initiating radical appears to bear a strong influence on metric studies [5-71.

483

Graft Copolymerization

the efficiency of grafting. Smets and coworkers [lo] found that when methylmethacrylate was polymerized in the presence of polystyrene initiated with benzoyl peroxide (BPO), an appreciable amount of grafting took place. When, however, the same reaction was initiated with azobisisobutyronitrile (AIBN) or ditertiarybutyl peroxide (DTBP), the amount of grafting was insignificant. Similar observationswere made by Misraet al. [ 113 during grafting of vinyl acetate onto starch in aqueous medium using BP0 and AIBN as radical initiators. Various methodsare available for effecting grafting. In most of them prior activation of the backbone polymer is involvedto afford active sites where vinyl monomers can be grafted. Different methodsof activation include: (1) physical activation, (2) chemical activation, and (3) radiation activation. A.

PhysicalActivation

Application of stress to polymeric backbonecauses segmental motions and molecularflow, which may lead to bond scission andconsequent formation of free radicals. Free radicals may also be generated by mastication and milling ofthe backbone polymeror can be accomplished by swelling it with suitable solvent. Penetration of the solvent through the polymer causes rupture of bonds leading to formation of free radicals. William and Stannett [l21 utilized freezing and thawing of the polymer monomermixturetechnique for effectinggrafting of vinyl monomers onto wool. Grinding of polyethylene wax with styrene in hexane gavea grafted product having goodflowability [ 131. Heat and impactresistant polymer moulding materials for automobiles were prepared by the injection mouldingprocedure for grafting unsaturated carboxylic acid monomers onto ethylene-propylene block copolymer [ 141. Thermosettingvinyl polymers graftedwithhydrolyzable silanes onto high-density polyethylene (HDPE) and vinyl trimethoxysilane-forming radicals on mechanicaltreatment [ 151. Production of a polymer with pendant basic moieties was performed by grafting dimethylaminoethyl-methacrylate to linear low-density polyethylene(LDPE) in the melt in a batch type internal mixer. The extent of grafting was determined by FTIR and nuclear magnetic resonance (NMR) spectroscopy, and the degree of crosslinking wasdetermined by measuring the product melt index.The grafted material was foundto be more stable suggesting an antioxidant effectof the grafted moieties[16]. Formation of graft copolymers hasalso been effected by carrying out the reaction at higher temperatures [17-191. Grafting in melt [20] or treating the polymer-monomer mixturein a mixer also produces graft copolymer [21,22]. CH3 B. ChemicalActivation Conventional radical initiators, various redox systems, P-diketonates of transition metals alone and in conjunc-

tion with additives, Lewis acids, and strong bases for chemical grafting have been used successfully for a long time. Graftingby use of any of the previously mentioned initiators proceeds with the initial formationof the free radicals, which initiate the polymerization of the vinyl monomer. Simultaneously,active sites are generated on the polymeric backbone where grafting cantake place. Grafting by free radical methods isdiscussed in the following sections. 1. RadicalInitiator In the presence of radical initiators such as benzoyl peroxide (BPO),azobisisobutyronitrile(AIBN), persulfates (S2OS2-), etc., grafting of vinyl monomers onto polymeric backbones involves generationof free radical sites by hydrogen abstraction and chaintransfer processes as described below:

I2 + 21'

1'+ P H + P

+ IH nm

l'+M+IM'vI-Mn-M' PH

+ I- Mn

-

+ M+PM'-P P' + I - M, -

P

M'+P

+I

- (M)n+1 - H

nm

- M, - M' M ' + P - (M)n+1 - I

(3)

where I2represents initiator molecule, PH is the polymer backbone and M is the vinyl monomer. The extent of grafting isdependent on various reaction parameters, such as concentrations of initiator and monomer, reaction time, temperature, etc. The nature of the initiator andthe monomer influencesthe reactivity of initiator and monomer toward grafting. It has been observed by Misra et al. that during grafting of vinyl monomers onto starch, wool, or rubber, conventional free radical initiators exhibit differentreactivity toward grafting.Whenvinyl acetate was grafted onto starch using BP0 and AIBNas the initiator, a lower percentage of grafting was obtained with AIBN[ 1 l]. The low reactivity of AIBNisexplained on the basis that the (CH3)+2*-C=N is resonance stabilized and less reactive. These radicals are, however, capable of initiating polymerization. CH3

CH3

CH3

I I CH3-C-N=N-C"CH3 I l b

I l

CH3"C'

e

C=N

N

-

CH3

I I

2CH3-C b

I

CH34

N

II

N

CH2H' vC

C=N'

I1

H34

I

b

N

(4)

Kaur and Misra

484

Misra et al. [l 11further observed that the molecular weight of poly(viny1 acetate) prepared by using either AIBN or BP0 was higher than the molecular weights of poly(viny1 acetate) isolated from starch-g-poly(viny1 acetate) graft copolymer andalso the molecular weights of poly(viny1 acetate) (PVAC) or isolated PVAC graft prepared by using AIBN as the initiator are higher than those prepared by using BPO.It is assumedthat benzoyl radicals, being highly reactive, terminate the growing chains with greater frequency and with consequent decrease in viscosity average molecular weight[ l l]. Acrylonitrile (AN) and methylacrylate (MA)havebeen grafted onto starch using BP0 as the initiator, and AN was found to be more reactive than MA. The observed reactivity was explained on the basis of polarity and solubility of the monomers in the reaction medium [23]. Modification of natural rubber has also been successfully attempted by grafting MA using BPO, potassium persulfate (KPS), and AIBN as radical initiators [24]. CH3

One molecule of acidfirst protonates TBHP to give

ond molecule of involving hydrogen bonding in the following manner:

0

I

0+H+

I

CH3 H

II

--s+H+S+H

II

0

(5)

When tertiary butyl hydrogenperoxide (TBHP) was used alone as the radical initiator, no graftingof methylmethacrylate (MMA) onto wool was observed. However, TBHP in conjunction with mineral acids, such as HzS04, HN03, or HC104 afforded good results [26]. Protonation of TBHP by the acid aided in the dissociation of TBHP to yield free radicals, which initiated grafting reaction.

I

a protonated complex (I) which , is stabilized by the sec-

I cH3BA>MA

CH3

I

CH3

Ethylacrylate (EA) and butylacrylate (BA) was grafted onto natural rubber using BP0 as the initiator [25]. KPS and AIBNwere ineffective in the production of the graft and the failure to form the graft isattributed to the “initiator effect.” The reactivity of different monomers towardgrafting onto natural rubber using different radicalinitiators was found to be:

H 0

I II II

0

(7)

The formation of the complex is expected to decrease the free energy of activation for the homolysis of the peroxide bond, and the decomposition of TBHP would occur at a lower temperature. It was further observed thatat a higher concentration of mineral acid, the decomposition of TBHP occurs via an ionic pathway, as reported by urne er [271.

Graft

Copolymerization

Different mineral acids showed the following reactivity order toward grafting in the TBHP-mineral acid system: (9)

&S04HNO3 > HC104

Chandel and Misra[28] observed that benzoyl peroxide alone was ineffective in the production of styrene grafted wool, but good grafting was obtained when used in conjunction with pyridine and acetic acid mixture, which acted as a pH modifier.The addition of pH modifiers along withpenetrating agents during BP0 initiated grafting of acrylic acid onto PET fiber was also studied by Suzuki et al. [29]. Persulfates have also been used as the radical initiatorfor graft copolymerizationof various vinyl monomers onto polymeric backbones. Grafting of

Polyethylene grafted with a wide variety of monomers using inorganic and organic peroxides shows improvement in thermal stability, elasticity, modulus, melt viscosity, and water resistance [31,32]. Carboxylic ion exchange resins [33] and chemical and heat resistant cation exchange membranes [341 have also been prepared bygraftingvinylmonomers onto polyethyleneusing chemical methods. It was observed by Borsig and coworkers [35] that in the presence of a reactive radical initiator such as

485

MMA onto cellulose wascarried out by Hecker de Carvalho and Alfred using ammoniumpotassium and persulfates as radical initiators [301. Radical initiators such as H202, BP0 dicumylperoxide, TBHP, etc. have also been used successfully for grafting vinyl monomers onto hydrocarbon backbones, such as polypropylene and polyethylene. The general mechanism seems to be that when the polymer is exposed to vinyl monomers in the presence of peroxide under conditions that permit decomposition of the peroxide to free radicals, the monomer becomes attached to the backbone of the polymer and pendantchains of vinyl monomersare grown on the active sites. The basic mechanism involves abstraction of a hydrogen from the polymer to form a free radical to which monomer adds:

dicumyl peroxide, polypropylene (PP)is not onlycrosslinked but also degraded. Cross-linking efficiency can be increased by polyfunctional monomers such as pentaerythritol diacrolein acetal (PEDAA) (a difunctional) or pentaerthritol tetraallyl ether (PETA) (a four functional monomer), and pentaerythritol difurfural acetal (PEDFA). The improved cross-linking efficiency of peroxides in the presence of polyfunctional monomers is due to the ability of the latter to add to radicals that are formed on the polymer chain via the double bonds of polyfunctional monomer.

+ or R - (CH = CH*), -~H"cH--cH~

I

CH3 2. Redox System In redox initiation, the free radicals that initiate the POlymerization are generated as transient intermediates in the course of redox reaction. Essentially this involves an electron transfer process followed by scissionto give free radicals. A wide variety of redox reactions, involving both organic and inorganic components either wholly

+ Polymer chain

or in part, may be used for this purpose. Because most of the redox initiators are ionic innature, their reactions take place in aqueous media and occur rather rapidly, even at relatively lowtemperatures. The general mechanism of grafting involvesan abstraction of the hydrogen atom from the backbone of the polymeric material by the transient radical formed during the redox reaction,

Kaur and Misra

486

thus generating macroradicalsonto which the monomer molecules add to produce graft copolymer. Commonly usedoxidants include peroxides, persulfates peroxydisulfates, permanganate, etc., and the reducing agents may be salts of metals such as Fe2+,C?+, V2+, Ti3+, Cu2+, oxyacids of sulfur, hydroxyacids, etc. There is no commonly accepted method of classifying redox initiators. They are subdivided in accordance with the nature of their main oxidizingcomponent. There exists a large number of redox systems that can initiate polymerization of vinyl monomers. The following are some of the few redox systems utilized effectively for graft copolymerization reactions. Peroxydisulfate System The activation of persulfates by various reductant viz. metals, oxidizable metals, metal complexes, salts of various oxyacidof sulfur, hydroxylamine, hydrazine, thiol, polyhydric phenols, etc. has been reported 136-381. Bertlett and Colman [39] investigated the effect of methanol onthe decomposition of persulfates and proposed the following mechanism. S2082- CH30H + HS04- + S04-0 + eH20H

+ SO4-* + CH30H + HS04- + eH2OH s2082-+ CH~OH+ HS04- + so4" + C H ~ O e H 2 0 H + e H 2 0 H + CH3OH + CH20 (12)

The Ag4 - K2S208 is also a very useful redox agent and the following mechanismfor its decomposition was proposed by Bawn and Margarison [40]. Ag+

+ S2082-+ Ag2+ + s04'- + Sod-*

+ S2OS2-+ Ag3+ + S04'- + S04-0 Ag3+ + 20H- + Ag+ + 26H (13) Ag2+

Polyamide fiber was grafted with MMA using the peroxydisulfate Ag+ redox system [411. Other metals such as Cu2+, Ti3+,and CO2+ have also been coupled with persulfate to form an effective redox system. A variety of vinyl monomers have been grafted onto different polymeric backbones using persulfate reducing agents as initiators. The K2S208 - NazS203--Cu ionaqueous or organic ligand complexsystem was used as the initiator for grafting methacrylic acidonto polyamide-6 fiber. It was observed that the system was most active when the fiber was treated with Cu-phenanthroline complex priorto polymerization. The copper complex in combination with Na2S203 and K2S2O8 in 25: 1 ratio fully inhibited polymerization. The initial rate of grafting was higher for unoriented fiber due to increased swelling, butthe yield of the graft copolymer was higher for oriented fiber due to inhibited chain termination [42]. Potassium persulfate with ferrous sulphate was successfully used as the initiator for grafting vinyl monomers onto wool [43] and cellulose [44] by Misra et al. They observed that there exists a critical molar ratio of

(oxidant)/(reductant)at which the percentage of grafting is at its maximum and beyond which it decreases. This was explained onthe basis that during Fe2+-induceddecomposition of SZOS'- excess of Fe2+consumes SO4-* to produce Fe3+ and S04'-. Fe3 at a higher concentration terminates the process. In order to establish the site of grafting on wool, graftingMA of and VAC was carried onto reduced wool. Reduction was effected by use of thioglycolic acid and the disulphide groups were converted to thiol (-SH) groups. The disulphide linkage is known to form a redox system with persulfate in the following manner. +

R-"-S--R + 03-S c! 0 so3+ RS+SO3+ &S + sod-*

(14)

Both fiS and S04-e directly or indirectly initiate polymerization. It was observed that the percentage of grafting ofMA onto reduced wool was higher than in wool. This suggested that the additional -SH groups provide sites for grafts [45]. R-SH

RS-

+ H+

Peroxydiphosphate System In conjunction with Ag+,V'+, CO2+,and acid, peroxydiphosphateforms an efficient redoxsystem for polymerization of vinyl monomers. H2P20g2- is assumed to be anactive species of peroxydiphosphate. The initiating species are OH and HP04-* and the termination is considered to be exclusively by mutual method. The following mechanism has been proposed for the redox reaction [46]. H2P20g2-

+ Ag+ + HP04" + HP04'- + Ag2+

HP04-* + Ag+ + HP04'-

+ Ag2+

+ HOH + 6 H + HzP04Ag2+ + HOH + 6 H + H + + Ag+ HP04-O + M -+fi-HPo4 (or GI) 6 H + M + &-OH(or GI) MI + nM + MI-M("-&I HP04-O

(16)

If the polymerization is carried out in the pl;esence of the backbone polymer, grafting occurs. Both OH and HP04"radicals can generate an active site on the backbone polymer by abstracting hydrogen, where grafting can take place. Peroxydiphosphate-V022+ redox system has also been investigated [47]. This system gave better results than when Ag+ was usedas the reductant. The rate of polymerization was found to increase with increasing temperature in the P2OS4- - VOZ2+system but not in the Ag+ containing redox system. Nayak et al. [48] graftedMMAontowoolusing perorydipbosphate in acidmediumandsuggested that H2PO4, OH, and

487


HP04-• are responsible for grafting. These radicals are capable of interacting with wool giving wool macroradicals where grafting can occur. Nayak et al. have also peroxydiphosphate-thiourea [49], successfully used peroxydiphosphate-Fe2+ [50], and peroxydisphosphate-Mn2+ [50] redox systems for grafting of vinyl monomers onto wool, silk, and polyamided fibers. Better results were obtained with the peroxydiphosphate-Mn2+ system. Ce4+Redox System Following the findings of Mino and Kaizerman [511 that ceric ion can form a redox system with cellulose, grafting onto various natural polymers has been carried out by the ceric ion method. In the case of cellulose, the reaction between ceric ion and cellulose occurs to produce active sites oncellulosein the following manner: Ce4+ Rce1,CH20H+ Complex + Ce3+, + H + + R,,~~CI~IOH (17) This is a highly selective process and very good results oncelluloseand starch graftinghavebeen observed. Ceric ion initiated grafting is usually carried out at lower temperatures and, therefore, wastage of monomer in chain transfer reactions is minimal. Extensive work on grafting of vinyl monomers onto differentpolymeric backbones hasbeen carried out using Ce4+ asthe redox initiator. Misra et al. have successfully carried out Ce4+ initiated graft copolymerization of vinyl monomers onto wool [52,53] and gelatin [54-561. In order to establish the site of grafting in wool, grafting has been attempted onto reduced wool and it was found that the " S H group of cystine of wool are involved in the grafting process [57,58]. Kinetic studies on ceric ion initiated graftingof AN onto sodium salt of partially carboxymethylated amylosehave been reported by Patel et al. [59]. When ceric ammonium nitrate (CAN) was used as a source of ceric ion, the presence of nitric acid was found to play a significant role. Ceric ion in water is believed to react in the following manner: Ce4+ H 2 0 + [Ce(0H),l3+ (i)

+

+

2[Ce(0H),l3+ "* [ C e " C e 1 6 +

+ Hz0

(ii)

(18) Thus, in aqueous medium, ceric ion exists as Ce4+ [Ce(0H),l3+ and [ C w e l 6 + . No grafting was observed in the absence of nitric acid. This is explained on the basis that in the absence of nitric acid, [ C e " C e I 6 + , because of large size, is unable to form a complexwithwool. However, with an increase in [HNO,], the equilibria [Eqs. 18 (i) and (ii)] shifttoward more and more of Ce4+ and [Ce(0H),l3+, which easily form complexes with functional groups of wool. The complex decomposesto give free radicals on wool where grafting occurs. At higher concentrations, these species adversely affect grafting efficiency.

Misra et al. have utilizedthe ceric-amine redox system for grafting MMA onto wool [601 and gelatin [61]. The graft yield was explained in terms of basicity, nucleophilicity, andsteric requirements of amines. A complex of ceric ion and amine(AH) decomposes to generate free radical species, which produce additional active sites onto the polymeric backbone where grafting can occur. WH + Ce4++Complex

I"*++

A - H + Ce4++Complex

A+wH+++AH

H + + Ce3+

+ H + + Ce3+ (19)

where WH = wool. When grafting of ethylacrylate was carried out onto gelatin using the ceric amine system [62], it was found that in the presence of more basic amines, which can form a complex withCe4+ with greater facility, a higher percentage of grafting was obtained. Besides basicity and nucleophilicity, steric requirements of the amines also playa part in the complex formation and these affect grafting. Similar results were also observed by Patel et al. [63] during grafting of AN onto sodium salt of partially carboxymethylated starch using a ceric-amine redox system. The effect of radioprotecting agents (RPA) on ceric ion-initiated grafting of MA onto cellulose has been studiedby Misra et al. [64]. Complex formation between the RPA and Ce4+ was assumed and it was foundthat threonine possessing " O H , -NH2, and "COOH groups can easily form a complex with Ce4+ and, therefore, increases percentage of grafting. Onthe other hand, in the presence of 5-hydroxy tryptophan and 5-hydroxy tryptamine, the percentage of grafting was found to be suppressed. This was explained onthe basis that the pheno1ic"OH groups present in these additives inhibit both polymerization and grafting. Complexformation between Ce4+ and the functional groups of the backbone polymer was also suggested by Chinese workers [65] during grafting of acrylamide onto 1,4-butanediol-4,4-diphenyl methane diisocyanate. Polytetramethyleneglycol copolymers capable of forming a complex between the carbamate groups and ceric ion lead to the generation of free radicals on the N-atom, which acts as the maximum site of grafting.

Metal Chelates Initiated Graft Copolymerization A number of metal chelates containing transition metals in their higher oxidation states areknown to decompose by one electron transfer process to generate free radical species, which mayinitiate graft copolymerization reactions. Different transition metals, such as Zn, Fe, V, CO,Cr, Al, etc., have been usedin the preparation of metal acetyl acetonates and other diketonates. Several studies demonstrated earlier that metal acetyl acetonates can be used as initiators for vinyl polymeriza-

488

Kaur and Misra

tion. Fewstudies have beenreported on the use of metal chelates as initiators for grafting. However, in recent years, attempts have been made by Misra et al. and Nayak et al. to effect grafting of different vinylmonomers onto starch, cellulose, and wool using various pdiketonates as initiators. Metal chelates are known to decompose upon heating to generate free radicals, which can abstract hydrogen atoms from the polymeric backbone producing an active site where grafting can take place.

"%

eH

/

"=% CH3

/CH3

+ PH-

"='\ + i, "=%

CH3

where PH is the polymer backbone, acac is the acetylacetone. Metal chelates afford a better initiating system as compared to other redox systems since the reactions can be carried out at low temperatures, thus avoiding wastage reactions due to chain transfer. Homopolymer formation is also minimum in these systems. It was observed by Misraet al. 166,671 that the maximum percentage of grafting occurs at a temperature much below the decompositiontemperature of the various metal chelates indicating that the chelate instead of undergoing spontaneousdecomposition receives some assistance either from the solvent or monomer or from bothfor the facile decomposition at lower temperature. The solvent or monomer assisted decomposition can be described as:

\C *;

\

CH3

2s (solvent)

I (acac),-

S

L

t S

+

(Y

+ e/=" \ CH3

Evidence for the formation of such a complex was obtained from IRspectroscopy [68] of the solid residue obtained by the treatment of MA with Fe(acac)3. Characterization of the complex by the ultraviolet (UV) spectral method also indicated the formation of the complex [68]. Misra et al. have carried out grafting of vinyl mono-

mers onto cellulose [68] and wool[69-74] using different metal chelates. The following reactivity order was established for grafting of MMA onto wool in the presence of various metal chelates as initiators. Fe(acac)3 > Cr(aca& - TBHP > M n ( a ~ a c ) ~ > Vo(acac)2


489

Nayak et al. have also reported grafting of vinyl monomers onto wool [75] and silk [761 using metal chelates as initiators. Hydrogen Peroxide-Reductant System Combinations of H202and Fe(I1) salts, generally known as “Fenton reagents,” have for many years been used in organic and polymer chemistry. Medalia and Kolthoff [77] summarized the principal features of this reaction. 1. In aqueous solution containing only Fe(I1) salts, H202,stable inorganic salts and acids, the stoichiometric reaction is followed if the iron salt is in excess over the peroxide.

H202

+ 2Fe2+ + 2H+ + 2Fe3+ + 2H20

(22)

2. When the peroxide is in large excess compared with the Fe(I1) less iron is oxidized than would correspond to the H202consumed if the above reaction was followed.The excess of the H202 is decomposed to oxygen and water. 3. In the presence of many organic compounds as well as of certain easily oxidized inorganic compounds such as the iodide ion, part of the oxidized H202is added to the compound.

the majority of the literature on graft copolymerization describes the use of a radiation method in some form or another. The growth and popularity of radiation as the initiating system stems from the improvement in the availability of ionizing radiationas a result of the introduction during recent years of more powerful nuclear reactors. Apart from beinginexpensive, it isa very convenient methodfor graft initiationas it allowsa considerable degree of control to be exercised over structural factors, such as the number and length of the grafted chains by careful selection of the dose and dose rate. Radiation is also unique in its ability to enable grafting to be carried out on prefabricated or shaped articles. The following sections describe, in a general sense, the variousmethods of radiationgrafting, their eficiency, applicability, andthe mechanism of the grafting reaction. Among the various methods used for effecting grafting by using high-energy radiations, the following methods deserve special mention: 1. The direct radiation graftingof a vinyl monomer

onto a polymer by the mutual method. 2. Grafting of vinyl monomers onto preirradiated polymers either in vacuum or air by the preirradiation method. 3. Grafting initiated by trapped radicals. 4. Intercrosslinking of two different polymers. 5. The doubleirradiationmethodinvolving the mutual method of grafting of vinyl monomers onto preirradiated polymer recently usedby Kaur et al. [85].

Following the discovery of Haber and Weiss [78] that Fe(I1) salts regct with H202 by one electron transfer process to give OH, Baxandale et al. [79] used the Fe2+ - H 2 0 2 system for effecting polymerization of vinyl monomers. Besides Fe2+, other reducing agents that may be Initiation by these methods can be carried out by used in conjunctionwithH202 are aliphatic amines, either gamma rays or by electron beam. Usually grafting Na2S203 thiourea, ascorbic acid, glyoxal, sulfuric acid, has been carried out either by the mutual or the preirraNaHS03, sodium nitrite, ferric nitrate, peroxidase, diation method usingCow as a source of gamma rays. AgN03, tartaric acid, hydroxylamine, ethylene sulfate, sodium phosphite, formic acid, ferrous ammonium sul1. Mutual Method phate, acetic acid, ferrous sulphate, and HN02, etc. In its simplest formthe direct grafting methodinvolves Fenton reagents have been used in numerous graft the irradiation of polymeric substrate in the absence or copolymerization reactions. Misra et al. have utilized presence of oxygen. Graft copolymerizationof the Fenton reagent as an initiatorfor grafting vinyl monomer monomer to the polymer is then initiated through the onto wool [80] and cellulose [81, 821 and observed that free radicals generated inthe latter. The reaction can be the molar ratio of [Fe2+]/[H2021influences percentage schematically written as: of grafting. Recently, Trivedi et al. [83] have grafted B AP acrylonitrile onto sodium Alginate using Fenton reagent as an initiator and observed that there exists a critical concentration of Fe2+and H202at which maximum percentage of grafting was obtained. They also studied the kinetics of reaction and foundthat the plots of lnRp versus MM], lnRp versus ln[H2021’”andlnRp versus In[Fe2+ bear a linear relation. This confirms the validity of the proposed reaction [84]. C.

Radiation-InducedGraft Copolymerization

Grafting by means of radiation is by far the most popular synthetic technique for modification of polymers, and

Kaur

490

where Ap W . and Ap&.Ap are polymeric free radicals derived from Ap, and R is a low-molecular weight radical or a hydrogen atom. Reaction (i) is expected to occur if Ap is a polymer of degrading type resulting in the formation of a block copolymer. If Ap crosslinks under irradiation, then reaction (ii)ismorelikely to occur leading to the formation of equal numbers of graft copolymer and homopolymer. addition, In the monomer B is also irradiated in the process to give the homopolymer: B - k A B q

(24)

In order to reduce the amount of homopolymer it may be advantageous to reduce the concentration of the monomer in the polymer monomer combination. Thus,

The hydroperoxide groups thus formed undergodecomposition duringirradiation producing macroxy radicals that offer sites for grafting.

-

-PA " H

-p*

+

6H (26)

(macroxy radical) The radiation sensitivity of a substrate is measured in terms of its GR value or free radical yield, which is the number of free radicals formed per 100 eV energy absorbed per gram. The highest grafting yields will occur for polymer monomer combinations in which the free radical yield of the polymer is much greater than for the monomer. It also follows that the grafting yieldwill increase at a lower monomer concentration. Graft copolymerization of different vinyl monomers has beensuccessfully attempted onto both natural [88-901 and synthetic polymers [91-933 by the mutual method. Percentages of grafting have been studied as a function of various reaction parameters that seem to influence the grafting reaction. Total dose used for grafting determines the number of grafting sites while the dose rate determines the length of the grafted branches. These can also be controlled by other factors such as presence of chain transfer agent, monomer concentration, reaction temperature, and amount of the solvent used. Misraet al. haveextensively studied the modification of natural polymers such as rayon [94],wool [95-971, and gelatin [98-991, as well as synthetic poly[100-102], isotacmers such as poly(viny1 alcohol) PVA

much higher yields are obtained if the polymer Ap is slightly swollen by the monomer B than ifAp is dissolved in a large excess of monomer B [861. Good results are obtained by swelling the polymer withthe monomer vapor [87]. A reduction in the concentration of monomer may further influence a number of other factors involved in the reactions such as viscosity of the reaction medium, the rate of chain propagation, the lengths of grafted branches, etc. If the grafting is carried out in air, the active sites on the polymeric backbone is attacked by atmospheric oxygen leadingto the formation of macroperoxy radical, which might abstract the hydrogen atom from the backbone polymer by an inter- or intramolecular process to give hydroperoxide groups as shown.

tic poiypropylene (IPP)[1031, and polyamide-6 [l041 by mutual method. In order to avoid chain transfer reaction, grafting was carried out in aqueous medium since water has zero chain transfer constant. It was observed that a critical valueof the total dose produce a maximum percentage of grafting beyond which the grafting percentage decreased. At a higher valueof total dose, bond scission or crosslinking of the backbone polymer may take place, which decreases the extent of grafting. A plausible mechanism of this reaction based on conventional oxidative degradation is outlined as follows:

In aqueous medium, excessive generation ofhydroxyl radicals from radiolysis of water preferentially initiate homopolymerizationof vinyl monomers leading to a lower percentage of grafting. Homopolymer, thus formed, if soluble in the reaction medium, increases the viscosity of the system, which results in the restriction of the mobility of the growing polymeric chains to the active site thereby decreasing the grafting level. In a medium of high viscosity, the primary radicals experience difficulty in addingto the active sites of the backbone polymer. With increasing total dose, recombination of primary radicals is accelerated leading to the formation of homopolymer at the expense of the graft. Another factor that affects the extent of grafting is the nature and the concentration of the monomer. Generally, maximum graft copolymer is observed at an opti-


491

mum monomer concentration beyondwhichgrafting percentage decreases. At higher concentrations, homopolymer formation becomes the preferred process. In addition, monomer chaintransfer reactions are also accelerated at higher concentrations. If the GR values of the monomer are higher as compared to that of the backbone polymer, homopolymer formation would abepreferred process. Use of excess of water decreased grafting percentage. This is because the excessive hydroxyl radicals producedas a result of radiolysis of water initiate polymerization at the expense of grafting. Also, if the monomer is soluble water, in an increased amount of water would entrap the monomer molecules and would decrease the rate and extent of grafting. The presence of alcohols in the aqueous medium generally decreases grafting. This is expected since the addition of alcohol breaks the tetrahedral hydrogen bonded structure of water and thus disturbs the association of active sites with water. This will lead to a decrease in grafting. In the presence of alcohols, chain

transfer reactions are accelerated leading to wastage of monomer and hence a decrease in grafting is observed. 2. PreirradiationMethod In this method the polymeric substrate (PH) is irradiated in air or oxygen to produce peroxide bonds [PH(02)] or hydroperoxide bonds[PH-OOH]. Irradiation of hydrocarbon usinga low-dose rate usually leads to the formation of hydroperoxide on the polymer backbone. Macrohydroperoxides upon heating decompose giving a hydroxyl radical, which can initiate the graft copolymerization of the added monomer.Polypropylenein particular appears to form a considerable amount of hydroperoxide groups on irradiationin air. The efficiency of the grafting reaction will depend directly on the kinetics of the radiation peroxidation process. The peroxide yield will necessarily depend on both the GR value of the backbone polymer andthe stability of the resulting peroxide at the irradiation temperature. The following processes have been suggested for peroxidation reactions:

I

AP A

0 -

-2

AP - 2

I!

nB

0.

Graft or block copolymers are expected to be formed viaprocess (28) (i) and (ii), respectively, depending on whether the polymer Ap isof a crosslinking type or a degrading type. No homopolymerization occurs in

2

1

OOH

-2

A

B this reaction other than by chain transfer to the monomer or by thermal initiation. For polymercontaininglabilehydrogen atoms, chain peroxidation can lead to hydroperoxide formation. The following reactions are expected to take place.

1 0 .

+6H” 2

1+ B

2 BqOH

Kaur and Misra

492

modification of various polymersby graft copolymerization using the mutual as well as the preirradiation methods [105-1091 and found that the preirradiation method produces better results [l 101.

Crosslinking and degradation reactions can predominate when the rate of diffusion of oxygen within the polymer is such that it cannot cope with the polymer radical formation. The concentration of polymer peroxides will increase with radiation dose. For low doses the peroxide concentration has been shown to build up linearly untilfor very high doses radiolysis of the peroxide groups becomes significant and reductionin the apparent overall rate of peroxidation is noticed. This may be the reason that at higher doses, the percentage of grafting was foundto decrease after attaining the maximum value at an optimum total dose. The advantage of this method over the mutual method is the avoidance of direct radiation polymerizationof the monomer and, hence, a higher percentage of grafting with minimum homopolymer formation obtained. is Misra et al. studied Intermolecular -CH&-

Another method has recently been utilized by Kaur et al. [85] involving grafting of a vinyl monomer onto preirradiatedbackbonepolymer by the mutualmethod. This method offers certain advantages over the conventional mutual method of grafting. Irradiation of the polymer in air prior to grafting introduces hydroperoxide groups in the backbone polymer by the intermolecular, intramolecular, or hydrogen abstraction process. CH3

CH3

l I H

3. DoubleIrradiationMethod

4 -CHZ"C-

I

L*

Intramolecular

Chien et al. [11 13 reported the formation of tertiary hydroperoxy groupsup to 10 adjacent carbon atoms during oxidative degradation. These hydroperoxide groups undergo radiolytic cleavageduringirradiation of the aqueous polymer monomer mixture.

gested that the hydroperoxide decomposition may take place by unimolecular or biomolecular reactions as follows:

-CH

JH3 +=

L

~6 + 6~

(Unimolecular)

H

(33)

CH3 (R = - C H A W )

Photolytic cleavage of hydroperoxy groups to give macroxy radical has been reported [ l 12-1 141. They sug-

I

CH3

-CHAW

+ H20 +

-CH2

Ao*

: 1

(34)

(Bimolecular) However, reaction (32) is expected to be more facile since it required minimum energy. Gray et al. [ 1151 have found that alkoxy radicals formed by the decomposition

of macrohydroperoxide groups can decompose by pscission reaction (35).

H3


493

CH3 CH3CH3H +

H H

-CHrd

It

+~

0

H T

I

H (35)

bOH

bOH Such a reciprocal motion of the kinetic chain (or back reaction) results in the decomposition of hydroperoxide groups withoutany interruption of kinetic chains and leads to the decrease in hydroperoxy groups. It has

Upon photolysis of .polypropylene hydroperoxide (PP-OOH) a major absorption at 1726 and 1718 cm" has been observed in the IR spectrum, which is attributed to the carbonyl groups. Sometimes the macroradical having free radical site reacts with a neighboring newlyborn hydroperoxide causing the formation of a macroalkoxy radical [l 161.

been suggested by many workers [117,118] that p-scission of tertiary alkoxy radicals is the major source of chain backbone scission during photodegradationof polypropylene.

The OH radical formed during the decomposition of hydroperoxy groups is very reactive and can either terminate a reaction (39) or may attack the weakbond of the backbone polymer resultingin the formation of a new radical (40).

CH3 """'CHr-(!-~Hz"''"'

+

6I-I

+

-CHr AH (39)

bOH On the basis of the above findings, graftingof vinyl monomers onto irradiated polypropylene has been attempted successfully bythe mutual method. Upon irradiation hydroperoxide groups are introduced, which provide sites for grafting. During mutual irradiation in the presence of the monomer inaqueous medium, these hydroperoxide groups and water undergo decomposi-

tion to yield macro hydroperoxy radical and hydroxyl radical. The hydroxyl radical initiates the polymerization of the monomer.

M

+ bH

"*

OH-M

OH

-

( M ) n - h (41)

The monomer canalso directly polymerize onto the macroperoxy radical as follows:

CH3 -CHd-CHp-

A*

+ nM

It was observed that beyond optimumtotal dose the percentage of grafting decreased. This may be due to the fact that at higher total doses beyond optimum, chain degradation by p-scission (reaction 35, 37, 38) occurs. Further at higher doses, hydroxyl radicals arising from

"+

*CHd AI(II1) > Zn(I1) complexes. D. PersulphateSystems

In aqueous solutions the persulphate ion is known as a strong oxidizing agent, either alone or with activators. Thus, it has been extensively used as the initiator of vinylpolymerization [43-471. However, only later, Kulkarni et al. [48] reported the graft copolymerization of AN onto cellulose using the Na2S203/K2S208 redox system. Gaylord et al. [49] reported the dilution and matrix effects in grafting of the styrene/AN binary mixtureonto cellulose with K2S208 as the initiator. Titledman and coworkers [501 reported the effect of hydroxypropylmethyl cellulose on the course of (NH4)2S2osdecomposition and claimeda route for grafting of vinyl monomers onto the polymer backbone. The decomposition of the peroxo salt, under the catalytic influence of the

Cell M + (n - 1) M --., Cell M, RM' + (m - 1) M + RM, Termination Cell M,, + Fe(II1) + Graft copolymer Cell M, + Cell M,, + Crosslinked graft copolymer RM', + Fe(II1) Homopolymer "+

Oxidation termination Cell' Cell'

+ Fe(II1) + Oxidation products + - S O i --., Dehydrocellulose + HS04-

E. Fenton's Reagent (Fe(lI)/H2O2) System

The use of Fenton's reagent (H202)as an effective system for grafting vinyl monomers onto preformed macromolecules was reported for the first time by Rogovin and coworkers [51-531. Rogovin [531 and Ogiwara et al. [54] suggested that the reaction of hydrogen peroxide and Fe2+ is a typical redox reaction. H M H

+ Fe(II)+

Fe(II1)

+ 'OH +

-OH

The OH radicals may initiate grafting viaabstraction of the hydrogen atom from cellulose or may also initiate homopolymerization. Grafting of methyl methacrylate on cellulose containing sulphonic acid groups [551and carboxy methyl groups [56] using the Fe2+/hydrogenperoxide redox system was investigated. Also, the rates of vinyl graftcopolymerization of AN and a-methyl styrene onto partially carboxymethylated cotton having different concentrations of carboxyl groupswere studied [57]. In this case,

Abdel-Bary and El-Nesr

506

the Fe2+-H202 redox system was used for initiation. The presence of sulphonic and carboxylic groups enables the iron ions to be in the vicinity of the cellulose backbone chain. In this case, the radicals formed can easily attack the cellulose chain leading to the formation of a cellulose macroradical. Grafting of methyl methacrylate on tertiary aminized cotton using the bi-sulphitehydrogen peroxide redox system was also investigated [W. The effect of Fe(I1) on grafting of 2-hydroxyethyl methacrylate onto polyester fibers in the presence of benzoyl peroxide was investigated [59]. It was found that increasing the iron ionconcentration decreases the graft yield. This suggest that excess Fe(I1) ions participate in the generation of free radical species and the iron ions seem to contribute to the termination and, consequently, decrease the graft yield. Thus, the reaction is not specific for initiation of grafting.Another disadvantage is that Fe(I1) ions formed-if not carefully removed-cause discoloration of the resulting product. The addition of Fe(I1)sulphoxylate is claimed to increase the rate of grafting, the yield of grafted polymer, and the conversion of monomer to polymer [60,61].The mechanism of grafting can be represented as follows:

+

Cell-H ‘OH + Cell’ + H20 Cell’ + M + Cell M OH M + H O M Cell M (n - 1) M + Cell M , HOM (m - l ) M + H O M m , + Fe(II1) + Graft copolymer Cell M M m Fe(II1) + Homopolymer Cell + ‘OH + Oxidation product

+

+ +

+

The reaction of any one of these two radicals with vinyl monomer leads to the formation of a grafted copolymer. IV. PHOTO-INDUCEDMETHOD The principal of photo-induced grafting the is absorption of electromagnetic radiation the in visible andultraviolet region. In this case, the molecule is said to bein an excited state. This requiresthat the polymer or the molecule should contain chromophoricgroups and should absorb the radiation in this region. The energy-rich molecule can either dissociate into reactive free radicals or dissipate this energy by fluorescence, phosphorescence, or collisional deactivation. The dissociation of polymer energy-rich moleculesinto reactive free radicals can be

F. Hydroperoxide and Peroxy Acid Redox

Systems

Peracetic acid formed in situ from acetic acid and hydrogen peroxide in acid aqueous solution was found to initiate graft copolymerization of vinyl monomers [62,63]. In the presence of Fe2+ and Fe3+ with peracetic acid, an optimum concentration was found to reach a high conversion of MMA, acrylamide, and acrylic acid, and a high yield of grafted cellulose [ a ] . The addition of a complexing agent, such as ethylenediamine tetraacetate (EDTA) to the Fe(I1) ions makes it possible to operate the grafting process at high pH. The following reaction mechanism was suggested:

+

AcOH + H202 + AcOOH H20 AcOOH + AcO’ + ‘OH AcOOH Fe2+ +-AcO’ + -OH + Fe3+ Cell H + AcO’ + Cell’ + AcOH , Cell’ nM + Cell M

+

+

In this case, two kinds of free radicals are formed leading to the formation of homopolymer and graft copolymer. The latter is due to the formation of cellulosic macroradicals. 111.

XANTHATEMETHOD

The xanthate method [62] is considered as one of the most promising methodsfor industrial chemical modification. The principal involved in the xanthate method of grafting is that cellulosic xanthate either ferrated or in acidic conditions reacts with hydrogen peroxide to produce macroradicals. The following reaction mechanism has been proposed:

used to initiate block and graft copolymerization. This process can be promoted bythe addition of photosensitizers. Such compounds may decompose easily into reactive free radicals or transfer their energy to other molecules in the system and, consequently, promote the graftcopolymerization reaction. The photosensitizers may exist freely in the reaction mixture or as pendant groups on the polymeric backbone. In the latter system they undergo facile homolysis to free radicals as reported by Norrish et al. [65, 661. The photo-induced process of modification of cellulose and its derivatives was reported by Geacintov and coworkers [67,68]. Thus, acrylonitrile, vinyl acetate, styrene, MMA, and the binary systemof styrene and AN were graftedonto cellulose and cellulosederivatives. In

Methods, Characterization, and Applicationsof Grafting

this case, photosensitive dyes, such as the disodium salt of anthraquinone (2,7-disulphonic acid), were used. Many reports are available onfree radical generation in cellulose and photochemical grafting of various monomers inthe presence of a large variety of initiators onto cellulose and cellulose derivatives [69-711. Photo-inducedgrafting onto woodcellulosewith several vinyl monomers using phenyl acetophenone and benzophenone derivatives as photosensitizers has been reported [72]. A fast increase in grafting was observed for short irradiation times, whereas longer ones resulted in a decreased percent grafting. Hon and Chan [73] examined the photo-induced free radical-generatedcelluloseand cellulose derivatives (methyl cellulose, ethyl cellulose, acetyl cellulose, hydroxyethyl cellulose, and carboxymethyl cellulose) using ESR spectroscopy. They reported that a high concentration of free radicals was generated in all samples 254 and irradiatedwithultraviolet of wavelength y 280 nm. Several types of free radicals were generated in cellulose due to chain scission, dehydrogenation, and decarboxymethylation reactions, whereas free radicals generated in cellulose derivatives were, by and large, due to the cleavage of the substituted side chains. Accordingly, grafting reactions of cellulose took place at the main backbone, whereas grafting reactions of cellulose derivativestook place at the substituted side chains. A higher degree of grafting is always achieved in a homogeneous medium than in a heterogeneous medium. The mechanism that explains the free radical generation on the cellulose derivatives can be represented as follows: Cel14CH3 hv_ C e l l 4 Cell"CHAH3

+ 'CH3

hv_ C e l l 4

+ 'CHAH3

The reaction between C e l l 4 radical and vinyl monomers leads to the formation of grafted cellulose. In the presence of photosensitizers generally used as photoinitiators, such as benzophenone and phenylacetophenone derivatives, the photoinitiator absorbs the UV radiation andtransforms to its singlet (S*) and then triplet (T*) states. After that it may decompose into free radicals or transfer its energy to cellulose or any other molecules in the system. Take benzophenone as an example: 0

II

PM-ph T*

4S*

_*

T*

+ CellOH "* p H + - P h + Cell' AH

Cell'

+ M + graft

507

If 2,2'-dimethyoxy-2-phenylacetophenone was used as the photoinitiator, the exited moleculedecomposes into a free radical as follows:

0 OMe

II

0

I

PHX-Ph

"*

S*

+ T* "*

II

P H '

AMe

OMe

+ 'CI-Ph

bMe

The formed radicals may abstract hydrogen from cellulose or a cellulose derivative leading to the formation of grafted cellulose or may initiate homopolymer formation. Grafting ofMMA onto dialdehyde cellulose proceeded more easily than onto mono-carbonyl cellulose [74]. In this case, the grafting reactions start with the photolysis of carbonyl groups. The mechanistic details of photo-induced graftingof styrene onto cellulose in the presence of mineral acids with methanol anddioxane as solvents was investigated [75]. It was foundthat acid enhances grafting and homopolymer formation. Analysis of homopolymers shows that acid reduces the chain length but increases the number of grafted chains. Akira and coworkers [76] reported photo-induced grafting onto cellulose in the presence of powerful oxidizing agentssuch as Ce(1V) and NaI04and free radicals generating aliphatic azo compounds (AIBN). The authors reported that grafting was initiatedby radicals via redox reactions between Ce(1V) andOH groups in cellulose, the grafting rate was affected by UV radiation. Photolysis of carbonyl groups formed as a result of oxidation in the presence of NaI04 enhances the grafting. In all cases the degree of grafting andthe apparent number of grafted chains increased linearly withan increase in the concentration of the chemical oxidants and the time of irradiation. Photo-induced homogeneous copolymerization of ethyl acrylate onto cellulose acetate using a series of photosensitizer combinations, such as aliphatic amines [77] and benzophenone and AIBN, benzoin and bromine-pyridine system [78], was studied. In addition, the kinetics of the grafting of ethyl cellulose with isobutyl methacrylate in homogeneous medium was studied using the conventional photosensitizers [79]. It was found that benzoin is more efficient than biacetal, benzil, and phenyl acetophenone, respectively.

V. GAMMA RADIATION-INDUCED GRAFTING The growth in popularity of radiation as the initiating system for grafting arises from the improvement in availability and cost of ionizing radiation. This is due to the introduction of more powerful nuclear reactors. Apart from its inexpensiveness, radiation isa very convenient method for graft initiationbecause it allows a considerable degreeof control to be exercised over such structural


508

factors as the number and length of the grafted chains by careful selectionof the dose and dose rate. Thus, the advantages of radiation-chemical methods are: (1) the ease of preparation as compared to conventional chemical methods, (2) the general applicabilityto a wide range of polymer combinations(due to the relatively unselective absorption of radiation in matter), and (3) the more efficient (and, thus, more economical) energy transfer provided by radiation compared to chemical methods requiring heat. A.

Mechanism of GammaRadiation Grafting

The theory of radiation-induced grafting has received extensive treatment. The direct effect of ionizing radiation in material isto produce active radical sites. A material’s sensitivity to radiation ionization is reflecteditsin G value, which represents the number of radicals in a specifictype (e.g., peroxyor allyl) produced in the material per 100 eV of energy absorbed. For example, the G value of poly(viny1 chloride) is10-15, of PE is 6-8, and of polystyrene is 1.5-3.Regarding monomers, the G value of methyl methacrylate is 11S , of acrylonitrile is 5.6, and of styrene is >0.69. During gamma radiation of polymer/monomer systems, numerous radiation chemicalprocesses involving macromolecules could, in principle, lead to the production of graft copolymers. The typical steps involved in free radical polymerization are also applicable to graft copolymerization includinginitiation, propagation, and chain transfer. However, the complicating roleof diffusion prevents any simple correlation of individual rate constants to the overall reaction rate. Among the various methods of radiation grafting,four have received special attention; these include: (1) the direct radiation of grafting of a vinyl monomer onto a polymer, (2) grafting on radiation-peroxidized polymer, (3) grafting initiated by trapped radicals, and (4) the intercrosslinkingof two different polymers.

B. DirectMethod In this method of grafting, the trunk polymer (Am) is irradiated while in contact with monomer B, existing either as a liquid, in a solution or vapor state. In this case, irradiation leads to the formation of active sites in the polymer (Am) or on the monomer (B). Polymer radicals either directly formed or via transfer reaction with the growing homopolymer radicals are directly used to initiate the graft copolymerization process. The resulting product is either a block or a graft copolymer depending onwhether Am is a degrading polymeror can be crosslinked under the effect of radiation [go]. This can be explained as follows:

A

A

A

A

A

B

A

A A A R is a low-molecular weight radicalor hydrogen atom. A and B are monomers. The homopolymer (B), arises from initiationby small radicalsR and also by radiolysis of the monomer B. B +. R

5R(B),

The rate of free radical production from Am and B are dependent on the G value of both: RiAm = G(Am) [I] [Am] Ri B = G(B) [I] [B] where Ri is the rate of initiation of the reaction, G(Am) and G(B)are the radiation chemical yieldsfor free radical production of polymer (Am) and monomer (B), respectively, and [I] is the dose rate. From the basic kineticsof radiation grafting,one of the most important properties in determining the efficiency of the process is the radiation sensitivity of the polymer, i.e., G(Am) relative to the monomer G(B) [81]. Thus, if G(Am) + G(B), grafting isfavored. The mutual technique works wellif the reaction yields a little homopolymer. A good example is the polyethylene/styrene system. If the contrary is the case, mainly homopolymer is formed andthe polymer substrate remains largelyunaffected. A good examplefor this case isthe PE/acrylic acid system where mainly polyacrylic acid is formed. However, methods have been worked out where a polymerization inhibitor is added to the monomer solution, thus greatly reducing the homopolymer formation [82]. If a solvent or swelling agent needs to be added, these should bechosen carefully, otherwise solvent radicals may predominate leading againto initiate homopolymerization. Dose effects are also important with the mutual method.If high-dose rates are used, the growing chains will rapidly terminate and graft yield and graft efficiency may result. For this reason, gamma sources such as gamma 6oCoor I3’Cs are preferred. Solvent or swelling agents are normally addedto increase monomer diffusion and enhance the efficiency and uniformity of the grafting. This can be achieved also with the vapor phase technique. However, nonsolvents have also been used to restrict the grafting close to the polymer substrate surfaces. This could be useful, for example, to improve adhesion or biocompatibility [83].


C. PreirradiationMethod In thismethod, the backbone polymer (Am) is irradiated in the absence of oxygen priorto exposure to the monomer (B), which may be ainvapor or liquid form. Irradiation of the polymer produces relatively stable trapped radicals. Graft copolymers resulted when the diffused monomers react with these active sites. The amount of formed homopolymers is very little by using this technique. However, the success of this method depends largely on the crystallinity of the substrate. Also, it depends onthe relative rates of reaction of monomers with trapped radicalsas well as the thermal decay of radicals at the higher temperature required for grafting. Thus, the graftingefficiency decreases withincreasing the decay rate of the radicals. Accordingly, it is advisable to conduct the grafting by irradiating the polymer at a low temperature and to allow the monomers to diffuse at the same temperature to prevent radical-radical combination on warming. If the polymer is semi-crystalline or below the glass transition temperature, the radical formed by radiation will remain trapped and able to react with the monomer to form a grafted side chain. One advantage of this method isthat an electron beam accelerator can be usedfor the irradiation step and the irradiated substrate can be reacted with the monomer in a continuous operation. This is valid for all vinyl monomers. One disadvantage of this method, is that, unlike the mutual technique, there is no protection of the polymer by the monomer or swelling agent. This can lead to degradation, crosslinking, or other changes in the polymer. Withmostgrafting systems, however, thishas not proved to be a major drawback. A second disadvantage is that the process is sensitive to the presence of air. This can be turnedinto a plus factor if the resulting peroxides are also used to initiate grafting [84]. D. PeroxidationMethod

Instead of using the polymer directly for grafting, it is possible to modify them by introducing a new function group via irradiation in air or oxygen. In this case, stable diperoxide or hydroperoxides are formed, depending on the nature of the polymer andthe irradiation conditions. These peroxy polymers canbe removed fromthe radiation source and stored at room or lower temperature until readyfor reaction. The reaction takes place in contact with monomer (B) in solvent and, if necessary, in air or vapor at an elevated temperature to form a graft copolymer. Catalysts can be added to accelerate free radical formation from the peroxy polymers and thus enhance grafting efficiency. The advantage of this technique is that the intermediate peroxy polymers may be kept for long periods of time before performing the final graftingstep. The grafting reactions using this technique can be described as follows:

509

a) Formation of diperoxide PH Po'

+

02

+ n(B)

P 4 0 - P a2Po' PMB),

-

if diperoxide formedat the chain ends, block copolymers are formed. b) formation of hydroperoxide PH + O2 y-rays, POOH 5P O P O nB __* PO(B), H 0 + n(B) HO(B), x + y = n

+

-

+ 'OH

Similarly, if the hydroperoxide is formedat theend group, a block copolymer is obtained. Experimental evidence is available to show that at room temperature diperoxides are formed in polyethylene, whereas polypropylene and polyvinyl chloride generate hydroperoxides. The temperature at which the peroxides decompose and initiate grafting depends on the type of polymer used. E. FactorsAffectingRadiationGrafting

There are a number of important factors that must be considered before applyinggammaradiation-induced grafting. These factors include the radiation sensitivity of the polymerhnonomer system, radiation dose and dose rate, type and concentration of inhibitor, type of solvent or diluent, and monomerconcentration. The effect of such parameters on the grafting efficiency during mutual grafting is given below. 1. Radiation Sensitivity of the Polymer1 Monomer System Ionizing radiation is unselective and has its effect on the monomer, the polymer, the solvent, and any other substances present in the system. The radiation sensitivity of a substrate is measured in terms of its G value or free radical yield G(R). Since radiation-induced grafting proceeds by generation of free radicals on the polymer as well as on the monomer, the highest graft yield is obtained when the free radical yield for the polymer is much greater than that for the monomer. Hence, the free radical yield plays animportant role in grafting process ~851.

2. Dose and Dose Rate Effects The dose and dose rate of irradiation are important factors inanyradiationgrafting system. In the direct method the total dose determines the number of grafting sites, while the dose rate determines the number of grafting sites, and the dose rate determines the length of the grafted branches. The length of the branches is also con-

510

trolled by other factors, such as the presence of chaintransfer reaction, the concentration of the monomer, the reaction temperature, the viscosity of the reacting medium, diffusionphenomena, etc. Diffusion of the monoAcrylamide merplays an important role in the direct radiation method. It is the way by which the monomer reaches the active sites within the polymer. It would beexpected that the rate of graft polymerization wouldAcrylamide be directly proportional to the radiation dose rate. In some cases, however, the diffusion of monomer cannot satisfy the increased rate of initiation within the polymer [86,87].

3. Effect of Temperature The effect of temperature on the kinetics of the direct radiation method isquite complex. Increase in temperature increases the monomer diffusion rate but also increases transfer and termination reaction rates of the growing chains and reduces the importance of the gel effect. Solubility and radical mobility may also change as the temperature is varied [88,89]. 4 . Effect of Inhibitors During mutual graft copolymerization, homopolymerization alwaysoccurs. This is one of the most important problems associated with this technique. Whenthis technique is applied to radiation-sensitive monomers such as acrylic acid, methacrylic acid, polyfunctional acrylates, and their esters, homopolymeris formed more rapidly than the graft. With the low-molecular weight acrylate esters, particularly ethyl acrylate, the homopolymerproblem is evidenced not so muchbyhigh yields as by erratic and irreproducible grafting. To avoid homopolymer formation, it is necessary to ensure true molecular contact between the monomer and the polymer. Even if this is initially established, it needs to be maintained during the radiation treatment while the monomer is undergoing conversion. Several methods are used for minimizing the homopolymer formation. These include the addition of metal cations, such as Cu(I1) and Fe(I1). However, by this metal ion technique, both grafting and homopolymerization are suppressed to a great extent, thus permitting reasonable yield of graft withlittle homopolymer contamination by the proper selectionof the optimum concentration of the inhibitor [83,90,91].

5. Effect of Solvent and Monomer Concentration The correct choice of solvents is essential to the success of radiation-induced graft copolymerization. Their influence on radiation graftinghas been the subject of many studies. It has been established that solvents play an important role in graftingbecause of the significance of polymer swelling. The grafting patterns to PP are solvent-dependent, thus, it is essential to examine the role of solvents. It was found that grafting in the alcohol is better than when other polar solvents such as dimethyl-

Abdel-Bary and El-Nesr Table 1

Substrate ReferencesMonomer PE Acrylic acid AN Styrene PP AN Acrylic acid Methacrylic acid Acrylates Styrene Styrene HEMA, NVP 4-Vinyl pyridine

PVC

%-98 99- 102 104, 105 92-94,103,105-109 110, 111 112 113 114

115

106,107 116 117 118

sulphoxide,dimethylformamide, acetone, or dioxane are used. The other important feature is the presence of an accelerated grafting effectin the alcohol, in dimethylsulphoxide, and dimethylformamide butnot in acetone and dioxane. The accelerating effect of methanol observed in the radiation grafting of styrene to polyethylene hasbeen attributed to the Thromsdorffeffect [92-941 similar to that found in the radiation graftingof styrene to cellulose [95]. In the case of cellulose, the wetting and swelling properties of the solvents as well as their radiation chemicalcharacterization were shown to be important in decreasing geleffects found with grafting of this trunk polymer. The first two properties are necessary to permit access of monomer to cellulose, since onlyhydrophilic solvents such as methanol, ethanol, and propanolwere found to be usefulfor grafting, whereas butanol was not. In Table 1a collection of some general polymer/ monomer systems in radiation grafting is given. These references are only representatives as the number of references in this area is very huge. VI. CHARACTERIZATION

A. Gravimetric Method The graft products are usually characterized by different methods. The first method isthe calculation of graft parameters known as the grafting percentage (GP), grafting efficiency (GE), andweight conversion percentage (WC). These parameters can be calculated as follows: A-B GP = -x 100 B A-B GE = Tx 100 L

WC%

=

A x B

100


where A,B, and C are theweights of the extracted graft product, substrate, andmonomer, respectively. This gravimetric method gives a direct and rapid indication about the graft reaction. Other characterization methods are usually used to detect the changes in physicalproperties, which usually result from the changes in the morphology and structures of the substrates due to grafting.

511

indicates there is some change the in crystallinity caused by grafting, comparison of the areas before and after grafting indicatesthis maybe a small effect. By assuming the difference in areas is due only to a difference in the amount of PE or PP present, in other words nodifference in the degree of crystallinity, the percent graft can be calculated from the equation:

B. Infrared where A, = area before grafting and A2 = area after Proof of grafting has generally been indicated by the grafting, p is the density. changes in the infrared spectra. The IR spectra of grafted cellulose has been examined by several investigators. D. ScanningElectronMicroscope This is valid for many other substrates. For example, the IR spectra of acrylonitrile graftedjute did not show The scanning electron microscope is generally used to a carbonyl peakthat was present in the physical mixture, detect the topography of the grafted surface, which usuindicating that grafting had occurred at this point [l 191. ally changesdue to grafted monomersonto thesurface. Kulkarni and Mehta[ 1201, while examiningthe IR specBesides, this method can be used also to detect the depth tra of acrylonitrile grafted chains after removal of the of grafting into the matrix. If binary monomer mixture cellulose backbone, found additional peak characteriswas usedfor grafting, the scanning electron micrographs tics of a " C O ester group indicatingthe presence helps to detect the grafted monomerdistributionby comof cellulose fragments. paring with micrographs of each grafted monomersepaThe amount of acrylonitrilegrafting onto PE powder rately. using an electron beam was determined from the absorbance of a nitrile group at 2240 cm" after extraction E. Swelling Measurements of homopolymer. In order to minimize the effects of a weighing error, an internal reference method utilizing Equilibrium swelling of grafted samples in the proper the methylene absorbance of PE at 730 cm" was solvent helps to detect the presence of grafted monomer. adopted. Thus, the mass of PAN in a sample wascorreFor instance, polyethylene does not swell practicallyin lated to the ratio of the absorbance A2240/A73,,, and the water. However, if polyethylene was grafted withwaterweight percent graft defined before was computed from soluble polymers suchas polyacrylic acid or polyacrylthe mass of PAN. amide, the equilibrium swellingof the product obtained Another technique often usedto verify actual graftmarkedly increases. Accordingly, the increase in swelling is to observe the change in the insolubility of the ing is an evidence of grafting. Incontrast, the ability of grafted cellulose in solvents such as cuene and cuprinatural rubber or styrene butadiene rubber vulcanizates ethylene diamine, which otherwise dissolve cellulose. to swell in gasolineor benzene markedly decreases due This change in solubility was found to be dependent to grafting with polyacrylonitrile. This decrease in swellupon the grafting level. In other solubility studies, where ing, again, is an evidence of grafting. the cellulose substrate and homopolymerwere both soluble, the grafted cellulose was completely dissolvedin F. Molecular Weight and Molecular Weight cuprammoniumsolution. The solutionwas acidified, precipitating the cellulose and grafted polymer and leavDistribution ing the homopolymer in solution. Molecular weight distribution of the grafted part is essential in order to precisely design the functional polyC. ThermalAnalysis meric membranes by application of radiation-induced graft polymerization andto control such a process. For In polymers having a certain degree of crystallinity, a example, the length and density of the polymer chains differential scanning calorimeter is used to determine grafted onto cellulose triacetate microfilteration memthe heat of fusion and, consequently, the changes in the brane will determine the permeability of the liquid degree of crystallinity inthe grafted and ungrafted samthrough and the adsorpitivity of the molecules on the ples. The changes in the crystallinity of PE found after functionalizedmicrofilterationmembrane. Thus, the grafting include a small, 2.5"C, drop in the location of molecular weight distribution of methyl methacrylate the maximum in the melting curve and a significant degrafted onto cellulose triacetate was carried out by the crease in the area under the melting peak 11211. Similar acid hydrolysis method of the substrate. From the gelresults were observed in case of grafting PP and PE/ permeation chromatogram, the molecular weightdistriEVA blends[981. While the decrease in the melting tembution was determined [122]. This method is valid only perature, represented by the shift in the melting curve,


512

when it is possible to degrade the substrate. In the case of grafted natural rubber, for example, ozonolysis is a very convenient process used to destroy the natural rubber segments while leavingthe plastomer chains intact [123].Alternatively, oxidation withperbenzoic acid can be used [ 1241.Osmometry or solution viscositymay then be usedto determine the molecular weightof the isolated nonrubber fraction.

G. DielectricRelaxation Dielectricrelaxation measurements of polyethylene graftedwith acrylic acid(AA), 2-hydroxyethyl methacrylate (HEMA) and their binary mixture were camed out in a trial to explore the molecular dynamics of the grafted samples [1251.Such measurements provide information about their molecular packing and interaction. It was possible to predict that the binary mixture used yields a random copolymer PE-g-P(AA/HEMA), which is greatly enriched with HEMA. This method of characterization is very interesting and is going to be developed in different polymer/monomer systems. VII.

APPLICATIONS

Since the changes in physical properties are often the impetus for grafting, it is necessary to briefly touch on this, in this section. A number of general reviews on graftinghave also included some discussion on the changes in physical properties [126-1291that usually determine the field of applications. Some other reviews deal with certain properties and applications, such as sorbency [70]and ion exchange properties [l301of cellulose. Grafting has often been used to change the moisture absorption andto transport properties of cellulosic materials when hydrophilic monomers, such as acrylamide, acrylic acid, and methacrylic acid, were grafted. Grafting of acrylic acid onto cellulose or starch can impart super water absorbing materialsfor use in sanitary napkins, diapers, and for soil stabilizer and other agricultural uses [ 131,1321. The dyeability of cellulose was found to be affected by grafting [ 1331 and has been examinedalongwith the corresponding lightfastnessand washfastness [l 191. Thus, for practical applications, continuous and semi-continuousmethods for grafting of partially carboxy methylcellulose were investigated [1341. Several authors have discussed the ion exchange potentials and membraneproperties of grafted cellulose [135,136].Radiationgrafting of anionicand cationic monomers to impart ion exchange properties to polymer filmsand other structures is rather promising. Thus, grafting of acrylamide and acrylic acid onto polyethylene, polyethylene/ethylene vinyl acetate copolymer as a blend [98],and waste rubber powder [137,1381,allows

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(1956). 125. A.A. Mansour, E. M. Abdel-Bary, E. M. El-Nesr, J . Elastomers Plastics, 26: 355 (1994). 126. R. B. Philips, R.B. Que're', J. G. Guiroy, and V. T. Stannett, Tappi, 55: 858 (1972). 127. J. C. Arthur, Jr.,Adv. Macromolec. Chem., 2 : 1 (1970). 128. J. C. Arthur, Jr., Adv. Chem. Ser., 9 9 : 321 (1971). 129. D. J. Mcdowall, B. S. Gupta, and V. T. Stannett, Prog. Polym. Sci., 10: 1 (1984). 130. R. Chatelin, Znf. Chem., 312: 181; C.A. 95: 99519x, 1981. 131. A. H. I. Zahran, J. L. Wiliams. and V. T. Stannett, J . Appl. Polym. Sci., 25: 535 (1980). 132. V. T. Stannett, G. F. Fanta, and W. M. Doane, Absor133. 134. 135. 136. 137.

138.

bency (P. K. Chattejee, ed.), Elsevier Science Inc., New York (1985). M. H. El-Rafe, E. M. Abdel-Bary, A. El-Hussini, and A. Hebeish, Angew. Makromol. Chem., 88: 193 (1980). A. Hebeish, M.H.El-Rafie, E. M. Abdel-Bary,and A. El-Hussini, Angew. Makrom. Chem., 88: 89 (1980). Y. Iwakura, T. Kurosaki,K.Uno,and Y. J. Imai, Polym. Sei., Part C4: 673 (1963). G. N. Richardsand E. F. T. White, J . Polym. Sci.,Part C4: 1251(1963). E. M. Abdel-Bary, M. A. Dessoki, E. M. El-Nesr, and M. Hassan, Proceedings3rd Arab International Conference on Polymer Science and Technology, Sept. 4-7, Mansoura, Egypt(1995); Po1ym.-Plastic Technol.Engineering, 36 (2), 241 (1997). G. Adam, A. Sebenik, U. Oserdkar, Z. Veksli, F. Ranogajec, and Veksij, Rubber Chemistry and Technology, 6 4 : 133(1991).

Applications Characterization, and Methods, 139.

of Grafting

H.Kaufer and E. M. Abdel-Bary, Colloid and Polymer

Sci., 260: 788 (1982). 140. A. K. Pikaev, L. V. Shemenkova, V. F. Timofeeva, and P. Y. Glazumov, Radiat. Phys. Chem.,35(13): 132 (1990). 141. A. A. Katbab, R. P. Burford, and J. L. Garnett, Radiat. Phys. Chem., 39(3): 293 (1992). 142. Haddadi-Asl, R. P. Burford, and J. L. Garnett, Radiat. Phys. Chem., 45(2): 191 (1995).

515 143. E. M. Abdel-Bary andE.

M. El-New,(unpublished work). 144. A. Nagaty, S . E. Shakra, S. T. Ibrahim, and 0 . Y.Mansour, Cell. Chem. Technol. 14: 177 (1980). 145. S. Kaizerman andG.Mino, Text. Res.J., 32: 136 (1962). 146. R. S. Rao and S. L. Kapur, J . Appl. Polym. Sci., 13: 2649 (1 969). 147. F. Ranogajec, M. Mlinac, I. Dvornik, Radiat.Phys. Chem. 18: 511 (1981).


Jayasree Konar, Bhola Nath Avasthi, and Ani1 K. Bhowmick Indian Institute of Technology, Kharagpur, India

1.

WHAT ARE SURFACE PROPERTIES?

techniquesprovide different, andsometimescomplementary, information. The most widely usedtechniques for surface analysis are Auger electron spectroscopy (AES), x-ray photoelectron spectroscopy (XPS), secondary ion mass spectroscopy (SIMS),Raman and infraredspectroscopy, and contact angle measurement. Some of these techniques have the ability to determine the composition of the outermostatomic layers, although each technique possesses its own special advantages and disadvantages.

Mostpolymers have verygoodbulk properties, but many industrial applications require them to have special surface properties as well. Surface treatments are used to change the chemical compositionof the surface, increase the surface energy, modify the crystalline morphology and surface topography, and remove contaminants and weak boundary layers. For example, good adhesion requires that the adhering surfaces be free of contaminants and that the substrate adequately wet the adherend. The removal of contaminantsfrom the surface is necessary to ensure good adhesion. Surfaces of adherends of low-surface energy, such as polymers, are treated to introduce certain functionally reactive groups necessary for the adhesive to be bonded. The surfaces may also be treated by the incorporation of a surfactant to change the surface tension of the adherend. The modification also helps to eliminate physical handling problems associated with static electricity and to enhance printing, decorating, wetting, and laminating qualities.

A.SpectroscopicTechniques I . Attenuated Total Reflectance Spectroscopy (ATR) An infrared beamis directed through a crystal of refractive index(n,) onto a sample of smaller refractive index (n2).The intensity of the reflected beam is monitored as afunction of the wavelength of the incident beam.These absorptions are used to identify the chemical structure. ATR has a sampling depth of about 0.3-3.0 microns.

II. MEASUREMENT OF SURFACE

2 . X-Ray Fluorescence Spectroscopy (XFS) This method allows us to characterize the first several hundred nanometers of depth of a solid. This may be an attachment to a scanning electron microscope (SEM). As the energetic electrons bombard the sample, ionization takes place. Ions with an electron vacancy in their atomic core rearrange to a lower energy state resulting in the release of electromagnetic energy of a specific wavelength. Analysisof the wavelengths of the X-radiation emitted identifies the atomic species present.

PROPERTIES

A number of techniques are available for determining the composition of a solid surface. Since the surface plays an important role in many processes, such as oxidation, discoloration, wear, and adhesion, these techniques have gained importance. The choice of a surface analysis technique depends upon such important considerations as sampling depth, surface information, analysis environment, and sample suitability. Different 517

Konar et al.

518

3. Auger Electron Spectroscopy (AES) This is a two-step process: an electron drops into a core vacancy and a secondary Auger electron escapes. The energy of the Auger electron depends on the chemical bonding state of the element from which itescaped. The maximum depth from which Auger electrons can escape isonly about 0.3-0.6 pm for mostmaterials. Thus, Auger spectroscopy is a technique that truly characterizes the surface region of the irradiated specimen. AES uses a low energy, 1-5 KeV, electron beam gunfor surface bombardment to minimize surface heating.

4. Electron Spectroscopy for Chemical Analysis (ESCA) or X-Ray Photoelectron Spectroscopy (XPS) In this method the surface is bombarded with low-energy x-rays, which are less disruptive than an electron beam. The energy is absorbed by ionization, resulting in the direct ejection of a core level electron, that is, a photoelectron. These electrons have an escape depth of less than a nanometer. Although X P S is less sensitive than AES, it provides a direct measure of the binding energy of core level electrons. and gives simpler spectral line shapes than AES. The following relationship holds: Energy of emitted photoelectron = Incident x-ray energy - Binding energy of ionized core-level electron

5. Zon-Scattering Spectroscopy ( I S S ) A solid sampleis bombarded with a stream of inert gas ions. Some of these ions are backscattered with some loss of energy after colliding with the surface atoms. Analysis of the scattered ion energies done is for identification of the surface atoms present.

Being sensitiveto thechemical composition the of outermost layers of the surface, contact angle measurement is widely usedfor characterizingpolymer surfaces. Surface characterization for polymers using contact angle measurement and X P S will now be described in detail, as these are the most widely used methods.

B. Contact AngleMeasurement Measurement of the contact angle at a solid-liquid interface is a widely used method for the determination of the surface energy of solid polymers. Fowkes [ l ] first proposed that the surface energy of a pure phase, ya, could berepresented by the sum of the contributionfrom different types of force components, especially the dispersion and the polar components, such that: where yad is the dispersion force component and yap is the polar force component. Fowkes [21 then suggested that the geometric meanof the dispersion force components was a reliable predictionof the interaction energies at the interface caused by dispersion forces. Thus, the interfacial free energy, yob, between phases a and b may be given by: where ya and yb are the surface free energies of phases

a and b, respectively. Owens and Wendt[3] and Kaelble

and Uy [4], by analogy withthe work by Fowkes, proposed another energy term, 2(yap ybP)'", to account for the effect of the polar forces (-yap and ybp are the polar components of the surface free energy of phases a and b, respectively). Thus,

6 . Secondary Zon Mass Spectroscopy (SIMS) The surface is bombarded with a stream of inert gas ions of energy, EO,and the sputtered target secondary ions of energy, E, are monitored, rather than the backscattered primary beam ions. Mass analysis of the secondary ions is camed out. The intensity and the energy are also determined. Each element has a characteristic value of E/ Eo. This allows the elemental analysis of the surface. Among the techniques mentioned previously, X P S has the greatest impact on polymer surface analysis. A major additional source of chemical information from polymers comes from IR and Raman spectroscopy methods. These vibrational data can be obtained from the bulk and the surface region, although the information depth is much greater than with AES, XPS, or ISS. A very simple, though indirect, method of surface analysis isthe measurement ofthe angle of contact that a liquid makes with the solid surface being analyzed. This method has been widely used to study changes introduced in a polymer surface by various treatments.

Considering a solid-liquid system, this relationship may be combined with the well-known Young'sequation to eliminate the interfacial free energy. Hence, cos0

= [-

1

+ 2(ysd yt)''z/y/ + 2(y,PYrp)'"/Y/] (4)

assuming the spreading pressure to be negligible and S and l to represent solid and liquid, respectively. The thermodynamic workof adhesion could also be obtained from the equilibrium contact angle 0 and the surface tension of the liquid as follows: Combining Eqs. (4) and (5): Various workers have used these equations extensively to understand the surface energetics of polymeric solids.

Surface Propertiesof Modified Polymers

519

I . ContactAngleHysteresis For a three phase system of solid-liquid-vapor in thermodynamic equilibrium, an angle O, is subtended at the triple line by the tangents onto the solid-liquid and liquid-vapor interfaces. In thermodynamic terms, this system parameter 0, depends, via Young’sequation, on the interfacial tensions. Instead of the equilibrium angle e, two other angles are obtained on real surfaces: the advancing angle 0, after enlarging the drop (wetting) and the receding angle0, after reducing the contact area of the drop with the solid (dewetting). The difference between the two angles is widely knownas hysteresis:

A0

=

0, - 0 ,

Both the advancing and receding angles be canobserved with one drop when the substrate is tilted [5]. The contact angle hysteresis of a polymer-liquid combination givesfurther insight intothe wetting behavior of a liquid onto a solid surface. Surface roughness and surface heterogeneity are the common causes of contact angle hysteresis. It is foundthat hysteresis becomes negligible when the roughness is 0.5-0.1 p m and when the heterogeneous phase is smaller than 0.1 pm.

2. Measurement of Contact Angles The most commonly used techniques for contact angle measurements are the sessile drop method and the Wilhelmy plate method. Results obtained from these two methods are in good agreement. 3. TangentMethod The tangent method is generally used for measuring the contact angle. It involves direct observation of the profile of a liquid drop resting on a plane solid surface. The contact angle is obtained by measuring the angle made between the tangent to the profile and the solid surface at their point of contact. This can bedone on a projected image or a photograph of the drop profile, or directly using a telescope fitted with a goniometer eyepiece. The polar and dispersioncomponents of the surface energy are generally obtained using twoliquids, for example water and formamide. To calculate ysp and y/, the following values for y/p and y t were taken [31: Water: y~ = 72.8, mN/m; yp ylP = 51.O mN/m Formamide: yr

= 58.2,

ylp = 18.7

C.

= 21.8,

mN/m; y p mN/m

mN/m;

= 39.5,

mN/m;

for polymeric systems. All elements, except hydrogen, are readily identifiableby XPS, since the different core level binding energies are highly characteristic. The concentration of different elements on a surface may be found by measuring the relative peak intensities and dividing them byappropriate sensitivity factors. Further, small shifts inthe binding energy ofa core level can be corroborated by considering the presence of different functional groups. For example, when a carbon atom is bonded to different groups of atoms of increasing electronegativity, a systematic shift inthe binding energy of the Cls peak isobserved: the higher the electronegativity of the group, the higher the binding energy of the Cls peak. 1. The XPS Experiment A basic XPS spectrometer consists of an x-ray source, an electron analyzer, a detector that counts the number of photoelectrons, and a data acquisition andprocessing system. XPS is conducted in a very high-vacuum environment (< torr). A high vacuumis also needed to prevent surface contamination and minimize scattering between the photoelectrons and gaseous molecules. A practical problem with using polymer in a high-vacuum environment is that low-molecular weightcomponents, additives, and water, may volatilize. In the majority of spectrometers, A1 and Mg are commonly used as x-ray target materials. With two anodes, AI and Mg, it is possible to resolve overlapping photoelectron and Auger electron peaks. This is because inanXPS spectrum the position of the Auger peaks changes if AI K, radiation is replaced by Mg K, radiation, but the positions of the photoelectron peaks are unaltered.

111.

SURFACEPROPERTIESOFMODIFIED POLYMERS

Over the past several decades, many surface modification methods for polymers have been developed. Pretreatment with corona discharge, gasplasma, flame, electron beam, and oxidizing solutions, and modification by blending, grafting, and copolymerization are common. Though many of these processes were developed primarily for treating polyethylene, they have been used with varying degrees of success for other polyolefinic materials, such as polypropylene, ethylene-propylene copolymers, and high-densitypolyethylene. The surface properties of polymers modified bythese techniques are now discussed briefly.

X-RayPhotoelectron Spectroscopy

Its ability to distinguish among different elements and different chemical bonding configurations has made XPS the most popularsurface analytical technique for providing structural, chemical bonding, and compositiondata

A.

ModificationbyChemicalMethods

This method is one of the oldest and most widely used techniques of improving the adhesion of polyolefins to metal and other substrates. Chemical treatment is usu-

2)

Konar et al.

520

allyused for irregularand particularly large articles when other treatment methods are not applicable. It involves immersion of the article in the etchant solution. Many types of chemical treatment are used in industry. Chromic,permanganic, sulphuric, andchlorosulphonic acids are often used as the oxidants. It has been shown that the adhesion of polyethylene to substrates, such as cellophane, steel, aluminium, and epoxy adhesives, improvesupon pretreatment withany of the etchants mentioned previously. Chromic acid etching has a significant effect upon surface topography of the etched polymer as shown by Blais et al.[6].They studied chromic-sulphuric acidetched low- and high-density polyethylene, as well as polypropylene, using transmissionelectron microscopy and found the etching rates to be independent of time upto 6 h at 70°C with the trend PP > LDPE >> HDPE. Reflection IR studies revealed extensive chemical changes on the surface in the case of LDPE but not HDPE or PP. New bondscorresponding to the introduction of -OH, > C 4 and-SO3H groups were detected. XPS studies of similarly treated LDPE and PP also confirmed both oxidation andsulphonation [7]. Yet another report on XPS studies revealed that the surface oxidation of PP rapidly reaches an equilibrium state and involves only a thin layer of the polymer (c=O (e.g., aldehyde or ketone) at 288.0 eV, a n d - C O w ( e . g . , carboxylicacid or ester) at 289.5 eV [21]. Some of these functional groups were also found inthe infrared spectrum of oxidized PP [14]. The 01s spectrum of the control sample shows a peak at about 531.5 eV, whereas the 0 1 s peak for oxidized PP is shifted slightly toward the high-binding energies (to about 533 eV). The shift of the 0 1 s peak toward high-

: 280292

‘m c

EPOXY present [161 ( W

Carbonyl present [ 171

286

0.016

528

534

540

528

534

540

@)

280

286

292

B i n d i n ge n e r g y

(CV)

Figure 1 CIS (left) and 01s (right) peaks for (a) PP and

(b) oxidized PP.

binding energy confirms that the relative concentration of carboxyl groupsincreases in oxidized PP.It is known that most oxygen functional groups give 01s binding energies at about 532 eV, whereas the ester oxygen in carboxyl groups shows a binding energy at about 533.5 eV [21].

Table 3 Binding Energies and Areas of Cls and 01s Spectra of PP and Oxidized PP

01s

Cls

Peak

Area Sample

position

PP 230.5 Oxidized PP

Area (eV/ms)

285.0 531.5

16.9

Peak (ev)

285.0 286.5 288.2 289.3

531 .S 254.4 (total) 165.8 248.4 65.2 6.5

533.7

20.8 255.4 (total) 7.0

Konar et al.

522

Table 4 Contact Angle and Contact Angle Hysteresis of Modified Polyethylene

of

Degree @a (H201(Hz01 grafting (mmoV100 Polymer (degree) (degree) g)(degree) (degree) (degree) (degree) LDPE

80 PEgVTMSl 78 PEgVTMSZ 67 PEgDBMl 64 PEgDBMz

PEgDBM3

73

20 25 13 17

20

12 13 1417

A@ WO)

o r

83

10

68 65 65 58 55 49 40 36

57

12 15

B. Modification by Graft Copolymerization Functionalization of polyolefinsthroughgraft copolymerization of unsaturated monomers containingpolar groups has received much attention in recent years. This type of functionalization is generally carried out for modification of some of the properties, such as adhesion and dyeability. Sometimes further crosslinking is done on these functionalized polymers. Graft copolymerization of an unsaturated polar monomeronto a preformed polymer is also widely recognized as a potential practical route to generate a novel class of interaction promoters (compatibilizers)in the case of multicomponent polymer systems. Vinyl silanes are among the most effective classes of such organofunctional monomers. They can be grafted easilyonto the backbone chains of polyolefins throughmeltprocessing. The technology ofmaking moisture-curable polyolefins through silane grafting by bulk processing has been developed [22,23]. The method of makingcrosslinkablepolyolefins through silane grafting has gained attention because of the various advantages, such as easy processing, low cost and capital investment, and favorable properties in the processed materials. Until recently, maleic anhydride has been used mostly for the functionalization of polyolefins due to the high reactivity of the anhydride group towardsucessive

0,

@a

A@

(HCONHz) (HCONH2) (HCONHz) 75 73 70 61 55 50

69

53

44

6 8 12 8 11

reactions [24-261. The maleic anhydride, although more reactive and effective than the corresponding ester, is very volatile, toxic, and corrosive. Hence, the ester dibutyl maleatehas been used ina few applications as it is more compatible with polyolefins than maleic anhydride W]. LDPE has been functionlized in bulk through dicumy1 peroxide (DCP)-initiated grafting of dibutyl maleate (DBM) and vinyltrimethoxy silane (VTMS)in the temperature range from140-200°C [28]. The advancing and receding contact angles of water andformamideon VTMS grafted PE (PEgVTMS) and DBM grafted PE (PEgDBM) are tabulated in Table4. The corresponding values of contact-angle hysteresis are also presented in that table. The advancing andthe receding anglesfor all the grafted PESdecrease when compared with the values , and 0 ,than of the control PE. PEgDBM shows lower 0 those of the silane system. These differences are more marked whenthe plane of PE is tiltedat 35". The change in O(H20) value is 12" for DBM-grafted PE1 as compared with PE. The contact-angle hysteresis value for grafted PES increases with an increase of the degree of grafting. It is observed that the total surface energy increases with grafting (Table5). The values become almost double for PEgDBM3. 'Theincreased surface energy of the grafted polymer comes mostly from the polar compo-

Table 5 Surface Energy of Grafted Polyethylene

of Polymer 12.69 13.35 13.63 15.36 15.46

LDPE PEgVTMSl PEgVTMSZ PEgDBMl PEgDBMz PEgDBM3

Degree grafting (mmo1/100 g )

Ysd

20 25 13 17 20

23.58 25.64 33.22 36.59 41.22

YSP

(mJ/m2)

(mJ/m2)

12.18

10.02 10.89 12.29 19.59 21.23 25.76

energy, Y (mJ/mZ) 22.20

Surface Properties of Modified Polymers

523

nent yg. y changes more than twofold for PEgDBM3. Owing to the presence of polar groups on the surface, the surface energies of grafted PES are expected to increase; the extent of this enhancement will depend on how many polar groups have been incorporated. The x-ray photoelectron (XP) spectra of pressed polyethylene indicates a low level of oxygenated species. The Cls and 01s regions of the XP spectra of PEgDBM and the Cls, Ols, and Si2p regions ofthe XP spectra of PEgVTMS are shown in Fig.2. The CISspectrum of polyethylene shows a single peak at 285.2 eV, indicative of the carbon-carbon bonds in PE. The Cls spectra of grafted polyethylene show newpeaks at highbinding energy (near 288-289 eV). The shift in the CIS peak indicates the presence of carbon-oxygen functionalities. To obtain atomic ratios, Ols, CIS,and Si2p peak areas have been corrected for different photoemission cross sections of these core electrons. These data are summarized in Table6. It is observed that the 0 :C ratio increases with grafting and, surprisingly, it is more for silane-grafted polymer. There is a four- to twelve-fold increase depending onthe level of grafting and DCP contents. An examination of Table 6 also confirms that increasing DCPconcentrationleads to increased 0 :C ratio for PEgVTMS. The increased 0 :C ratio for PEgDBM might be due to the presence ofDBM groups on the polymer backbone. The appearance of a new peak at the higher binding energy region (288-289 eV) of the Cls

c 1s

Table 6 X P S Data of Grafted Polyethylene

Polymer

Degree of grafting (mrnoU100 g)

-

LDPE PEgVTMSl PEgVTMS2 PEgDBMl PEgDBM2 PEgDBM3

% Si to

% 0 to

total c

total C

-

0.65 6.95 7.52 2.30 4.12 6.90

20 25 13 17 20

4.2 6.5

-

spectra can be ascribed to the carboxylic ester group 1211. The higher 0 :C and Si :C ratios for PEgVTMS can be explainedby considering the incorporation of the -Si0-C group onto the polymer backbone during grafting. At higher values of the 0 :C ratio, lowering of the contact angle isexpected. Higher 0 :C ratio means more functional groups are present and, naturally, one can expect an enhanced surface energy of the solid. Loweringof the contact angle suggests better wetting and, therefore, enhancement of surface energy. However, it is interesting to note here that the surface energy of PEgVTMS is lower, although the 0 : C ratio is higher than that of PEgDBM. It is well knownthat the silicones

Si 2p

0 1s

x 30

x 30 I*

534 532 530 289 287 285 283

I

~71 i I 1 L

98

100

102

104

B i n d i n g energy ( C V ) Figure 2 Binding energies (in eV) of Cls, Ols,and Si2p peaks from surfaces of (a) LDPE, (b) PEgDBM, and (c) PEgVTMS.

524

Konar et al.

are release agents and have very low surface energy [29]. Apparently, the expected increase in ys due to increased 0 :C ratio is partially nullified by the presence of silicon atoms.

C. Modification by ElectronBeam Treatment

Electron beam-initiated modification of polymers is a relatively new technique with certain advantages over conventional processes. Absence of catalyst residue, complete control of the temperature, a solvent-free system, and a source of an enormous amount of radicals and ions are some of the reasons why this technique has gained commercialimportance in recent years. The modification of polyethylene (PE) for heat-shrinkable products using thistechnique has been recently reported [30,31]. Such modification is expected to alter the surface properties of PE and leadto improved adhesion and dyeability. The equilibrium contact angles for PE and triallyl cyanurate (TAC) grafted PE with water and formamide are presented in Table 7. The grafting of TAC onto PE increases the surface energies of modified PES [32]. Table 7 reveals that the grafting of TAC onto PE decreases the equilibrium contact angles of water and formamide from 92" to 65" and from 75" to 53", respectively. Thisdecrease is a function of the monomer level and the irradiation dose. At a fixed irradiation dose of 15 Mrad, variation of the TAC level from0.5 to 3 parts causes a reduction inthe contact angles of water by 13" (from 88" to 75") and of formamide by 11" (from 72" to 61"). This is due to the fact that the concentration of

polar groupson the surface increases as a result of grafting of the polyfunctional unsaturated monomer. These polar groups are evident from the IR and XPS studies [32]. With increasing irradiationdose at a constant TAC level of 1 part, the contact angles of water and formamide are lowered by7" up to an irradiation dose of 10 Mrad, andthen increased for up to a 20 Mrad irradiation dose. The results can be explained interms of the level of surface modification described later. The effect of irradiation of samples in the absence of TAC on the contact angles is also reported in Table 7. Modification of the surface takes place, as is evident from the decrease in the contact angles of water and formamide. The change, which is maximum at anirradiation dose of 10 Mrad, is due to the generation of polar functionalities on the surface. This is also corroborated from the IR/XPS studies described later. The contact angles are lowered further when TAC isincorporated in the system (compare TO/S with T1/5, TOM with T1/15, etc.) The results of surface energy of grafted PES are reported in Table 7. It is observed that the dispersion component of the surface energy, ysddoes not increase with the incorporation of TAC, althoughthere is an approximately two- to three-fold increase in the polar component of the surface energy, resulting in an overall increase of the surface energy value.A change inthe total surface energy of the samples without TACdue to irradiation (i.e., T0/2, T0/5, T0/10, etc.) is also evident. With the increase in irradiation dose, however, there is an optimum value of the surface energy at 10 Mrad irradiation. In order to explain these results, the grafting levels calculated from the IR spectra [301 are

Table 7 Contact Angle 0 and Surface Energies of PE and TAC Grafted PES

level (mmoV100 Sample

72

72

TO10 TO.5/15 T1.5/15 T2/15 T3/15 T1/2 TU5 TV10 T1115 T 1/20 TO12 T0/5 TO110 T0/15 TO120

Grafting O (mJ/m*) (mJ/m2) g) (mJ/m2) (degrees) (degrees)

4.8 5.1 15.0 25.0 5.7 6.7 7.4 4.6 2.6

20 20 20 19 18

-

18

-

16 16

-

92 88 82 79 75 69 65 70 88 82 78 80 84

O(HCONH2) YSP 75 72 67 64

61 60 58 53 60 62 73 72 68 69 69

TI115 for example, means, sample with 1 part TAC irradiated at 15 Mrad.

Ys

YSd

24 26 32 39

21 20 20 17 17 15

28 26

19

4 6 8 10 12 15 17 19 17 16 6 10 12 11 7

28 31 34 35 34 33 24 25 27


525

given in Table 7. It is clear that the changes in surface energy values are generally in accord with the changes in the grafting level. Asthe grafting levelincreases, the surface energy values increase. IR studies of irradiated samples without TAC indicate absorbance peaks at 1140 and 1732 cm" due to the generation of " C H r and >C=O functionalities. The ratio of the peak at 1732 cm" with respect to that at 1470 cm" was calculated for all samples. On introduction of TAC, the ratio increases by four- to eightfold depending on the amount of TAC andthe irradiation dose. Figure 3 shows the high-resolution spectra for PE in the Ols, Nls,and Cls regions and illustrates several points of interest. Three representative samples, that is, control PE and grafted PES at 15 Mrad irradiation dose using 2 parts TAC (T2/15) and at 10 Mrad irradiation dose with 1 part TAC (TIAO), are shown in the figure. The Cls spectrum for the control PE film surface shows a peak at 284.5eV indicative of the carbon bonds in PE with a small shoulder, probably related to the presence of an 01s peak. This control sample shows a small percentage of oxygen but no nitrogen on the surface before grafting. After grafting, several percentages of oxygen and a small percentage of nitrogen are incorporated into the PE surface, and the Cls spectrum shows new peaks toward higher binding energy, indicating the formation of carbon-oxygen and carbon-nitrogen functionalities.A simple deconvolution of the spectrum for Cls in Fig. 3c is shown. The high-binding energy regionof the CIS spectrum can be fitted with three peaks corresponding to carbon atoms with a single bond to oxygen or nitrogen at 287.0 eV, carbon atoms with two bonds to oxygen at 288.0 eV, and carbon atoms with three bonds to oxygen at 290.0 eV. The 01s spectrum of the control sample shows a peak at about 532.2 eV, while the 01s peak for grafted PES shifts slightly toward higher binding energies (532.9 eV) (Table 8). This shift confirms that the relative concentration of carboxyl and carbonylgroups increases in grafted PE. Also, the 01s peak area increases three times on modification.The grafted samples show a Nls

(a)

-2 c

280 285 290 Binding m r g y (CV)

395 WO 405 5275 532'5 5375 Bindiq energy W ) Bindiq mwgy(eV)

(b) Binding cncqy (CV)

Binding energy (CV) Bindiq -(*V)

280 285 290 (c) Binding energy (@V

395 400 405 5275 5325 5325 Binding energy (*V) Binding energy (*V)

Figure 3 (a) Core level spectra of polyethylene (TO/O). (b) Core levelspectra of polyethylene graftedwith 2 parts TAC at an irradiation dose of 15 Mrad (T2/15). (c) Cls, Nls, and 01s peaks for polyethylene grafted with 1 part TAC at an irradiation dose of 10 Mrad (TlAO).

Table 8 X P S Details of Cls, Nls, and 01s Spectra from TAC Grafted PE

Cls Peak center

PeakArea

center

Sample TO/O T2/15 20.31 TV10 20.35 TO/10 20.28 T0/15

285.5 284.9 285.0

01s

Nls

-

W15 for example, means, sample with 2 parts TAC irradiated at 15 Mrad.

Peak

Area

2.02

-

center

Area

532.2 532.4 532.6

7.30 7.41 2.23 2.64

526

Konar et al.

peak at 400.7 eV, and the incorporation of a nitrogen peak suggestsa grafting reaction. It must be pointedout here that the samples even without TAC showa surface concentration of oxygen on irradiation in line with the IR studies. The O/C ratio, however, increases significantly for the TAC grafted PE. D. Modification by Blending

Multiphase polymers, such as block and graft copolymers and polymer blends, show distinct chemical and physical properties at the surface [33,34]. The low surface energy component generally comes to the surface because the thermodynamic driving force reduces the total free energy of the system. Siloxane polymers are commonly used in copolymers or blends to reduce surface energy andthereby reduce their coefficient of friction. The surfaces of most blockcopolymers and blends of siloxane polymers are rich in siloxane [35]. Similar behavior is apparent in the blend of equal parts by weight of silicone and EPDM, where silicone diffuses onto the surface during aging [36]. Surface segregation of these polymers occurs only at verylow concentrations. Hence, this method has been recently proposed to modify the surface without affecting the bulk properties of polymers. The equilibrium contact angles andsurface energies of silicone, EPDM, anda 50:50 blend of Si: EPDM are 9. On aging,at 175"C, the contact angle reported in Table of water gradually decreases with aging time and then increases at 48 h of aging. For all the rubbers, the surface energy increases with an increase of time of aging inthe initial stage and then decreases. The peak value after which there is a decrease in the surface energy is 9 h for the silicone and the blends and 24 h for the EPDM. In the blend of silicone and EPDM (50:50 by weight) the surface energy values in the initial stages of aging are similar. Butthese are lowered for the blends at longer times of aging. It is clear that the polar contribution increases from 7.95 mJ/m2for the unaged sampleof EPDM to 50.55 mJ/m2 for a 24 h aged sample, a six-fold increase; from 1.86 mJ/m2 for the unaged sample of silicone rubber to 16.40 mJ/m2for the 9 h aged sample, a

nine-fold increase; and from8.73 mJ/m2for the blend to 17.19 mJ/m2 for the 9 h aged sample. These values decrease on continued aging. The dispersion component, on the other hand, shows a decrease in general before the final increase. There is a clear indication of increased polar functionalities, mainly due to >C=O function, from IR measurements. For example, during aging at 175°C for 48 h, there is about a 25% increase of >C* functionalities for the control EPDM, while a 15% increase of the same group is registered for the EPDM in the blend. Clearly, the EPDM in the blend is protected from oxidation by the presence of silicone rubber. From the XPS measurements of 50:50 silicone-EPDM blend, the concentrations of carbon, oxygen, and silicone have been calculated [36]. It is observed that the relative area percent, which is proportional to the concentration of a particular chemical species, changes due to oxidation. The total peak area of 01s increases due to oxidation (Fig.4). Silicone peaks Si2p and Si2s appear at about 102 and 153.5 eV, respectively, in the blend. It is interesting to note that the peak area of silicone increases from 31.74 to 46.6 units (47% increase) on aging for 9 h at 175°C. It can be inferred that the silicone diffusesonto the surface during aging. The total concentration of silicone is 18.5% after aging, as compared with 12.4% on the unaged surface. Concomittantly, there is approximately a 7% increase in the oxygen concentration. The Cls concentration, however, is lowered by about 13%. These results are in general agreement withthe contact angle values and IR observations [36]. E. Modification by Corona Treatment

Thistechniqueisdesigned to generate a sufficiently high-voltage electrical discharge at the surface of the moving substrate, which may be in the form of sheet or film. The corona treatment functions at atmospheric pressure and relatively high temperature. It works with both electrically conducting and insulating substrates. There is substantial roughening of the surface following corona discharge treatment [37]. Very significant surface oxidation also occurs [38]. XPS techniques identify

Table 9 Contact Angle of Water and Surface Energies of EPDM, Silicone and Their50: 50 Blend

(mJ/m2) ys energy) (surface

O(H2O) (degrees) Status of sample EPDM Silicone Unaged 23 Air aging at 175°C For 3 h 94 For 9 h 19.30 102 For 24 h 104 107 For 48 h 100

50 :50EPDM Silicone blend 95 87 80 7018.96 67 86

105

50: 50 blend 15.47.OO

90 15.74

18.73

90 36.08

11.41

90

-11.34

25.08 51.22

15.47


H

.-

. .

291 287 283

531

527

535

E 3

.2

534 530 287 283

1 0 0 156 152 106

53L 530 287 283

1 0 0155 151 104

W E

x 333

(a) 534 530 287 283

l00 B i n d i n ge n e r g y

l04

W x 3'33

155 151

(,V)

Figure 4 Cls, Ols, Si2s, and Si2p peaks of: (a) unaged 50: 50 silicone-EPDM blend; (b) 50: 50 silicone-EPDM blend aged for 9 h at 175°C;(c) silicone rubber aged for 9 h at 175°C;and (d) EPDM aged for 9 h at 175°C.

the presence of hydroxyl, ether, ester, hydroperoxide, aldehyde, carbonyl, or carboxylic groups in corona-discharge treated polyolefins. Briggs andKendall[39] have provided some insight into the adhesive interaction between discharge-treated LDPE and a commercial nitrocellulose based printing ink. Using reagents that block specific functional groups, they showed that if enolic OH was eliminated fromthe discharge treated surface, then ink adhesion droppedto zero. On the other hand, elimination of carboxylic acid functions brought about only a slight loss of ink adhesion. By analogy with autohesion findings, it was concluded that hydrogen binding between enolicOH groups andan acceptor moiety of the printing ink was mainly responsible for the adhesion. F.

Modification by Gas Plasma Treatment

Gas plasmatreatment operates at low pressure and relatively low temperature. While the corona treatment is applicable to substrates in sheet or film form, the gas plasma process can treat objects of virtually any shape. The gases most widely usedto generate plasma byfreeradical reactions include air, argon, helium, nitrogen, and oxygen. All these, with the exception of oxygen,

behave as inert gases. Oxygen plasma is normally used to clean surfaces by oxidizing them, resulting in the formation of -C&, " O H , or U groups. Argon, helium, and nitrogen plasmas improve surface wettability andbondabilitybymeans of surface activation and crosslinking. The diffusion of electrons plays a major role in the surface treatment of polymers in a gas plasma apparatus. The adhesion of PE tapes increases to a maximum by this treatment at 0.20 to 0.30 8, for 20 to 30 S [40]. IR absorption at 1600-1750 cm" increases 1.5-fold on discharge treatment owing to the formation of polar >C=O and "COOH groups, which increase tape adhesion. Brosse et al. [41] modified isotactic polypropylene isotactic polyand other polyolefins bya cold plasma. In propylene, plasmatreatment results in a polypropylene crystallization of paracrystalline or smectic form into a a-crystalline form. Further, the active films are susceptible to react with monomers in a postgrafting reaction. Tetrafluoromethaneplasma modifiedpolyethylenes indicate two mechanisms: degradation and fluorination. These reactions are competitive and parallel [42]. G.

Modification by Flame Treatment

Flames are also plasmas, characterized by electron densities of about 108/cm3 and electron energies of about 0.5 eV. Many excited species are present in the flame, namely free radicals, ions, excited atoms and molecules, and electrons [43]. Excited species that have been observed include 0, OH, NH, NO, and CH [M]. Flame treatment is predominantly used with articles of relatively thicksection, such as blow mouldedbottles, although ithas been applied to polyolefin films as well. The most important variables in the process are the air-gas ratio and their rate of flow, the nature of the gas, the separation between burner and surface, and the exposure time.

H. Modification by Other Methods UV irradiation on a polymer surface produces chemical modification as well as wettability and bondability improvement. It causes chain scission and oxidation on polymer surfaces \even in the presence of an inert gas [45]. Carbonyls are found to be introduced onto polyethylenes on UV irradiation. Sivram et al. [46] have used photochemical treatments for surface modification of polymers. They have generated surfaces of vaying surface energies by simple organicreactions. Thermal aging is another simple pretreatment process that can effectively improveadhesion properties of polymers. Polyethylene becomes wettable and bondable by exposing to a blast of hot (-SOOOC) air [47]. Meltextruded polyethylene gets oxidized andas a result, carbonyl, carboxyl, and hydroperoxide groups are introduced onto the surface [48].

Konar et al.

The effect of thermal aging onpolyethylene and isotactic polypropylene have been studiedby Konar et al. [49]. They used contact angle, contact angle hysteresis, and XPS to characterize the modified surfaces of the polymers. Hysteresis increased with agingtemperature. In the case of polyethylene, thermal aging led to a significant increase in adhesion strength of polyethylene with aluminium, butthe increase in the case of polypropylene was much less marked. Oxidative degradation of nitrile and hydrogenated nitrile rubber was studied using IR, X P S , and contact angle measurement [50]. The contact angle of rubbers decreased with aging time. XPS studies indicated that the oxidation of hydrogenated nitrilerubber takes place through-CkN, whereas the double bondsare attacked in nitrile rubber. Surface modification by plasma, laser and ozone treatment has been discussed in a recent book Dl].

REFERENCES 1. F. M. Fowkes, J . Phys. Chem., 67: 2538 (1963). 2. F. M. Fowkes, Treatise on Adhesion and Adhesives (R.

L. Patrick, ed.), Marcel Dekker, Inc., New York, Vol. 1, p. 344, (1967). 3. D. K. Owens and R. C. Wendt, J . Appl. Polym. Sci., 13:

1740 (1969). 4. D. H. Kaelble and K. C. Uy, J . Adhesion, 2: 50 (1970). 5. S . Wu, Polymer Interface and Adhesion, Marcel Dekker, Inc., New York p. 15, (1982). 6. P. Blais, D.J. Carlsson, G . W. Csullog, and D. M. Wiles, J . Colloid Interface Sci., 47: 636 (1974). 7. H. A. Willis and V. J. I. Zichy, Polymer Surfaces (D. T. 8.

9. 10. 11.

12. 13. 14. 15. 16. 17.

Clark and W. J. Feast, eds.), Wiley Interscience,. New York p. 287, (1978). D. Briggs, D. M. Brewis, and M. B. Konieczko, J . Mat. Sci., 11: 1270(1976). A. Baszkin, L. Ter-Minassian-Saraga, and C.R. Lisbeth, Acad. Sci., Paris, Ser C , 268: 315 (1969). C. Foresca, J. M. Perena, J. G. Faton, and A. Bello, J . Mat. Sci., 20: 3283 (1985). D. W. Dwight and W. M. Riggs,J . Colloid Interface Sci., 4 7 650 (1974). E. H. Andrews and A.J. Kinloch, Proc. R. Soc. Lond., A332: 385 (1973). J. Konar, S . Ghosh, A.K. Banthia, and R. Ghosh, J . Appl. Polym. Sci., 34: 431 (1987). J . Konar and P. Maity, J . Mat. Sci. Letts.13: 197 (1994). J. Konar, G . Samanta, B. N. Avasthi, and A.K. Sen, Polym. Degradation Stability, 43: 209 (1994). J. Konar and R. Ghosh, Polym. Degradation Stability, 21: 263(1988). J. Konar and R. Ghosh, J . Appl. Polym. Sci., 40: 719

(1%). 18. J. Konar, R. Ghosh, and S. K. Ghosh, Polym. Degradation Stability, 22: 43 (1988). 19. J. Konar and R. Ghosh, J . Adhesion Sci. Technol., 3: 609 (1989). 20. J. Konar, R. Ghosh, and A. K. Banthia, Polym. Comm m . , 29: 36 (1988).

21. D. Briggs, V. J. I. Zichy, D. M. Brewis. J. Comyn, R. H. Dahm, M. A. Green, and M. B. Konieczko, Surf. Interface Anal., 2: 107 (1980). 22. H. G . Scott, U.S. Patent 3,646,155 (Feb. 29, 1972). 23. P. Swarbrick, W. J. Green, and C.Maillefer, U.S. Patent 4,117,195 (Sep. 26, 1978). 24. N. G . Gaylord and M. Mehta. J . Polym. Sci. Polym. Lett., 20: 481 (1982). 25. N. G . Gaylord, M. Mehta, and R. Mehta, J . Appl. Polym. Sci., 33: 2549 (1987). 26. Y. Minoura, M. Ueda, S . Mizunuma, and M. Oba, J . Appl. Polym. Sci., 13: 1625 (1969). 27. R. Greco, G . Maglio and P. V. Musto, J . Appl. Polym. Sci., 33: 2513 (1987). 28. J. Konar, A.K. Sen, and A.K. Bhowmick, J . Appl. Polym. Sci., 48: 1579 (1993). 29. K. E. Polmanteer, Handbook of Elastomers-New De-

velopments and Technology (A. K. Bhowmick and H. L. Stephens, eds.), Marcel Dekker, Inc., New York,p. 551,

(1988). 30. T. K. Chaki, S. Roy, A. B. Majali, R. S. Despande, V. K. Tikku, and A. K. Bhowmick, J . Appl. Polym. Sci., 53: 141(1994). 31. T. K. Chaki, A. B. Majali, R. S . Despande, V. K. Tikku,

and A. K. Bhowmick,

61 (1994).

Angew. Makromol. Chem., 217:

32. J. Konar and Ani1 K.Bhowmick, J . Adhesion Sci. Technol., 8: 1169(1994). 33. D. W. Fakes, M. C.Davies, A. Brown, andJ. M. Newton, Surf. Interface Sci., 13: 233 (1988). 34. H. Inoue, A. Matsumoto, K. Matsukawa, A. Ueda, and S . Nagai, J . Appl. Polym. Sci., 41: 1815 (1990). 35. N. M. Patel, D. W. Dwight, J. L. Hedrick, D. C. Webster, and J. E. McGrath, Macromolecules, 21: 2689 (1988). 36. Ani1 K. Bhowmick, J. Konar, S. Kole, and S . Narayanan, J . Appl. Polym. Sci., 57: 631 (1995). 37. K. Rossman, J . Polym. Sci., 19: 141 (1956). 38. C. Y. Kim and D. A. I. Goring, J . Appl. Polym. Sci., 15: 1357 (1971). 39. D. Briggs and C. R. Kendall, Surf. Interface Anal., I : 189 (1980). 40. V. I . Bukhgatler, R. I. Belova, N. V. Evdokimova, and A. L. Goldenberg, Plast. Massy., 2: 56 (1981). 41. F. Poncin-Epaillard, B.Chevet, andJ. C. Brosse, J . Appl. Sci.. 53: 1291(1994). 42. F. Poncin-Epaillard, B. Pomepui, and J. C. Brosse, J . Polym. Sci. Polym. Chem., Pt.A . 31: 2671 (1993). 43. J. R. Hollahan and A. T. Bell, Techniques and Applications of PlasmaChemistry, Wiley Interscience,New York (1974).

44. A. G. Gaydon, The Spectroscopy of Flames, Chapman

and Hall, London (1974). 45. N. J. DeLolis, Rubber Chem. Technol., 46: 549 (1973). 46. S . Sivram, Lecture at the Indo-French Symposium at P&M Curie University, Paris (1995). 47. W. H. Kreidl and F. Hartmann, Plast. Technol., I : 31 (1955). Brewis, and M.B. Konieczko, Eur. 48. D.Briggs,D.M. Polym. J., 14: 1(1978). 49. J. Konar, A. K.Bhowmick, and M. L. Mukhejee, J . Surface Sci. Technol., 8: 331 (1992). 50. S. Bhattacharya, Ani1 K. Bhowmick, and B. N. Avasthi, Polym. Degradation Stability, 31: 71 (1991). 51. K. L. Mittal, Polymer Surface Modification: Relevance to adhesion, VSP, Utrecht, Netherlands, 1996.

34 Grafting of Cellulose Abd-Alla M. A. Nada National Research Centre, Cairo, Egypt Mohamed Adel Yousef Helwan University, Cairo, Egypt

Most of the grafting methods involve the creation of free radicals on the cellulose moleculeto initiate the graft polymerization. Free radicals are produced by yradiation [3,4]or by chemical means [5]. The most important way of producing radicals and initiating grafting is the utilization of redox systems [6]. Hydrogen peroxide andferrous salts have been usedto initiate polymerization [7].The possibility to use thiocarbonate (xanthate method) for the preparation of graft copolymers has been recognized by Faessinger andConta [SI.A graft copolymer isformed, whose composition and properties depend on the type of monomer and on the conditions of the reaction. Briekman [9]stated that the xanthate grafting process does not require a special atmosphere. For optimum efficiency of the grafting, the presence of some solvents is usually necessary.

I. INTRODUCTION

Grafting copolymerization of different monomers onto cellulose is clarified.Several oxidant systems have been used to generate free radicals on cellulose molecules to initiate graft copolymerizationas photoactive quinones, ozone, hydroxyl radicals, ceric ammonium nitrate, potassium bromate, andpotassiumperminganate. Partially, xanthate cellulose isalso used for the preparation of grafting. On the other hand, the effects of different physical treatments, e.g., beating, grinding, and cellulose swelling withethylene diamine, onthe reactivity of cellulose toward graftingare studied. Also the effect of partial carboxymethylationand acetylation with a different degree of substitution (DS) on the grafting of cellulose has beenstudied. Finally, the effect of grafting medium (aqueous or solvent) on the cellulose graftability with monomer is discussed.

111.

II. GRAFTCOPOLYMERIZATION OF CELLULOSE

Grafting presents a means of modifying the cellulose molecule through the creation of branches of synthetic polymers, which impartto the cellulose certain desirable properties without destroying the properties of cellulose. The polymerization of vinyl monomersmay be initiated by free radicals or by certain ions. Depending on the monomer, one or the other type of initiation may be preferred. The grafting process depends on the reactivity of the monomer used, the type of initiation, and cellulose accessibility [l ,2]. 529

INITIATION BY FREERADICALS

Free radicals can be generated on the cellulose chain by hydrogen abstraction, oxidation, the ceric ion method, diazotization, introduction of unsaturated groups, or by y-irradiation. If a vinyl monomer is polymerized in the presence of cellulose by a free radical process, a hydrogen atom may beabstracted from the cellulose by a growing chain radical (chain transfer) or by a radical formed by the polymerization catalyst (initiator). This leaves an unshared electron on the cellulose chain that is capable of initiating grafting. As cellulose is a very poor transfer agent [IO],very little copolymer results from the abstraction of hydrogen atoms by a growing chain radical.The

Nada and Yousef

530

only way in which a chain-transfer reaction between a growing chain radical andcellulose can be made to succeed is by first introducing into the cellulose molecule certain atoms, or groups, that are readily abstracted by radicals. Therefore, in most cases a radical produced by the initiator is responsible for the formation of the graft copolymer [l 1,121. This technique is based in the fact that when cellulose is oxidized by ceric salts such as ceric ammonium nitrate {Ce(NH4)2(NO&}free radicals capable of initiating vinyl polymerization are formed on the cellulose. However, the possibility remainsthat the radical formed is an oxygen radical or that the radical is formed on the C-2 or C-3 instead of the C-6 carbon atom. Another mechanism, proposed by Livshits and coworkers [13], involves the oxidation of the glycolic portion of the anhydroglucose unit. Several workers [14,15], however, have found evidence for the formation of some homopolymer. In the ceric ion method free radicals are first generated and are then capable of initiating the grafting process [16-181.

A. Grafting of Cellulose with Ceric Ions Grafting of pulpwith methylmethacrylate monomer using a ceric salt redox system is carried out. The effect of different variablesfor example, monomer dose, reaction time,reaction temperature, acid concentration, initiator concentration, and liquor ratio/is studied. The effect of the presence of different amounts of residual lignin in pulp on the grafting process is demonstrated. Residual ligninpercentage in the pulp plays an important role in determining the grafting rate.

l . Effect of Monomer Concentration Figure 1 illustrates the effect of monomer concentration (0.5-3 mL/g pulp) at 25°C for 1 h, 0.1% initiator, and

1% acid concentration at 30: 1 liquor ratio on the total conversionandgrafting efficiency percentage. Total monomer conversion percent is increased from 34% to 126% by increasing the monomer concentration from 0.5 to 3 mWg cellulose. Onthe other hand, it isevident that grafting efficiencyincreases as the methylmethacrylate (MMA) concentration increases from 0.5 to 1 m.L. Enhancement of the grafting efficiency due to theincrease of the monomer concentration can be interpreted in terms ofimmobilityof the cellulose macroradicals. Hence, availability of MMA molecules in the proximity of these radicals was a prerequisite for graft initiation and propagation.The increase in grafting efficiencycan also be attributed to the fact that the increase of monomer concentration from 0.5 to 1 mL/g pulp allows the graft formation to grow, and long graft chains are formed since the active sites are the same in all monomer concentrations. By increasing the monomer concentration more than 1 mL, the grafting efficiency is decreased, since this increase allows the formation of more homo-

401

2ot 0

x .

.

.

.

1m1 monomer 2 / g 3pulp

Figure 1 Effect of monomer concentration on the total conversion percent and grafting efficiency.LR 30: 1, acid concentration l%, initiatorconcentration O.l%, grafting time 1 h, and reaction temperature27°C 0-0 = total conversion (%); o"0 = grafting efficiency (%).

polymers. It seems that both grafting and polymerization are favored at high-monomer concentration, but homopolymerization prevailsover grafting.

2. Effect of Reaction Time Total monomer conversion percent is increased by increasing the reaction time witha slow rate, reaching the maximum value (60.6%) after a reaction time of 1 h. Increasing the reaction time more than 1 h causes the reaction rate to nearly level off. The leveling off can be ascribed to the depletion in monomer and initiator as well as to the shortage of available grafting sites as the reaction proceeds. However, the grafting efficiencyhas the same trendas the total monomer conversion percent; that is, it is increased by increasing the reaction time and reaches its maximum value after 1 h. By increasing the graftingtimestill further, graftingefficiencyis slightly decreased or nearly leveled off, indicating that the formation of homopolymer predominates. The maximum grafting efficiency(59.6%) is obtained after a reaction time of 1 h.

3. Effect of Acid Concentration The effect of acid concentration on the rate of grafting has beenstudied. It is clear that increasing the acid concentration from 0.5% to 1.0% increases not onlythe total monomer conversion percent from 60.1% to 61.5%, but also the grafting efficiency from 58.0% to 59.6%. This

Grafting of Cellulose

531

can be explained as follows: at 1% acid concentration (H+) is sufficient to establish a balance between the suppression of the rate of active species formation and the rate of primary radicalgeneration. By increasing the acid concentration more than l%, the total monomer conversionandgraftingefficiency are decreased. Increase of acid concentration over 1% may cause oxidation of the formed free radicals andthereby suppress the rate of initiation.

4. Effect of Liquor Ratio The effect of changing the grafting liquor ratio on the total monomer conversion and grafting efficiency ranging from 20: 1 to 60: 1 was studied. Total monomer conversion percent as well as grafting efficiency increased is by increasing the liquor ratio from 20: 1 to 40: 1. Increasing the liquor ratio up to 40: 1 allows more homogeneous distribution of the chemicals andincreases the mobility of the monomer, which,in turn, increases its penetration velocity onto cellulose macroradicals and, consequently, increases the grafting efficiency.

5 . EfSect of Initiator Concentration Figure 2 shows the effect of ceric ammonium nitrate concentration on the total monomer conversion and grafting efficiency of the produced grafted pulp. From this figure, it is clear that the total conversion percent increases with an increasing initiator concentration, reaching a maximum valueof 64.9% using a concentration of 0.15%, and grafting efficiency reaches its maximum value (59.6%)at 0.1% initiator concentration. This increase in grafting efficiencycan be discussed in terms

Ft

of the increasing rate of the termination step as ceric ion increases the redox process, which generates more radical sites on the cellulose, and then propagates by the addition of monomer.

6 . EfSect of Reaction Temperature Total conversion percent is increased by increasing the reaction temperature from 10°C to 50°C. A sharp increase is observed up to 30"C, then it is increased slightly by increasing the temperature up to 50°C. On the other hand, the grafting efficiency increases with increasing reaction temperature and reaches its maximum value (66.9%)at 30°C. It is well knownthat the rising reaction temperature, up to 30"C, enhances the grafting efficiency through one or more of the following factors: (1) increased mobility of the monomer and, hence, a higher rate of monomer diffusion from the reaction medium into cellulose; (2) reaction of the already formed homopolymer chain with cellulose macroradicals; (3) increased cellulose swellability; and (4) enhancement in the rate of initiation and propagationof the grafting process. Increasing the grafting temperature more than 30°C decreases the grafting efficiency.

7 . Effect of the Residual Lignin in Pulp Graft polymerization of methylmethacrylate monomer onto pulps of different residual lignin contents using ceric ammonium nitrate as the initiator was carried out to studythe influence of this residual ligninon the graftability of these pulps (Fig. 3). From this figure one can

loo

8o

%

40

2o 0

L 1

2

initiator

0 2

3

4

6

8

Residual lignin

10

12

%

concentration

Figure 2 Effect of initiator concentrator on total conver-

sion percent and grafting efficiency. LR 30: 1, reaction time 1 h, reaction temperature 27"C, monomer concentration 1 mL/g pulp, and acid concentration 1% W = total conversion (%);

20

o"0 = grafting efficiency (%).

Figure 3 Effect of residual lignin percent in the pulp on total conversion percent and graft yield percent. LR 40: l , reaction time1 h, reaction temperature 27"C, monomer concentration 1 mL/g pulp, initiator concentration l%, and acid conversion 1%; W = total conversion (%), 0-0 = graft yield (%).

Nada andYousef

532

notice that decreasing the residual lignin percent in the pulp enhances not only the total conversion percent but also the graft yield percent. A sharp decrease in total conversion and graft yield percent is noticed upon increasing the residual lignin percent in pulp from 1.68% to 4.97%. The graft yield and total conversion percent reach nearly a zero value at 12% residual lignin in the grafted pulp. This can be attributed to the fact that increasing the residual lignin in pulp increases the consumption of ceric ions due to the oxidizing effectof ceric ammonium nitrate on the polyhydric alcohol of lignin; this results in the formation of quinonoid structure and, hence, significantly retards and/or inhibits the grafting reaction, which lowers the propagation of the reaction [19]. Moreover, the presence of lignin, which acts as a cement to the cellulosefibers, decreases the penetration of the chemicals throughthe cellulose fibers, and, consequently, decreases the grafting reaction. To sumup, the optimum conditionsfor methylmethacrylate grafting onto pulp by the ceric ion redox system can be summarized as follows: the grafting is done at 30°C for a l-h reaction time, using liquorratio 40: 1, acid concentration l%, initiator concentration O.l%,and monomer 1 mL/g pulp. Free radicals can also be formed on the cellulose molecule by using H202 [20-231, ozone-oxygen mixture [21,22], perminganate [22], and bromate [24]. B. Grafting by Xanthate Method [25]

Cotton linters and viscose grade wood pulp were partially xanthated under different conditions to study the effect of the degree of substitution on the acrylamide grafting of these pulps. Sodium hydroxide solutions of 2%, 4%, and6%were used andthe vapor phase xanthation process was applied for 0.5, 1.0, 1.5, and 2.0 h for

i:;

each concentration. Grafting of the partially xanthated cellulose samples was carried out under the same conditions. The results obtained show that there is no direct relation betweenthe degree of substitution and the graftingyield. The most important factors that affect the grafting process by the xanthate method are sodium hydroxide concentration and time of xanthation. Cotton linters and viscose grade wood pulp were grafted with acrylamide using the xanthate method. The effects of monomer concentration, reaction time, hydrogen peroxide concentration, reaction temperature, and liquor ratio on the grafting process were studied. Optimum conditions for the grafting reaction were established [26].

+ C = S += Cell - OH + CS2 I I

Cell-OH

+ Monomer += Cell - OH

S

0 '

b

A

0' .

b Monomer

0

I . Effect of Monomer Concentration The effect of acrylamide monomerconcentration on the grafting percentage, grafting efficiency, polymer loading, and polymerization efficiencyare shown in Fig. 4. Increasing the monomer concentration from 0.5 to 3 g/ g cellulose increases the grafting from 21.3% to 59.5% and from 12.8% to 37.5% for cotton linters and wood pulp, respectively. On the other hand, the grafting efficiency decreases by increasing the monomer concentration. The grafting and polymerization efficiencyof the grafted cotton linters were higher than those of wood pulp. This can be explained by the differences in the chemical and physicalstructures of the two pulpsrather than the degree of substitution of the partially xanthated pulps.

50

c3 ).( 20

I

Cell-OH

t

-0

- 0

1 2 3 monomer z pulp

1 2 3 monomer : pulp

Figure 4 Effect of monomer: pulp ratio on grafting yield and grafting efficiency. 0 = cotton linters; 0 = wood pulp.

533


attributed to the exothermic character of the polymeriza2. Effect of Grafting Time Polymer loading, polymerization efficiency, and graftingtion reaction; also increasing reaction temperature may cause the oxidation of some of the created free radicals. of cotton linters and viscose wood pulp were increased Therefore, both the grafting and polymer loading were with increasing the grafting time. Inthe first 30 min the decreased by increasing the grafting temperature. reaction was slow, then its rate increased until 60 min; after 60 min the reaction rate acquires a leveling value. C. Grafting of Cellulose by Using This can be ascribed to the depletion of the monomer y-Radiation and initiator concentration as well as a shortage of the available grafting sites as the reaction proceeds. The Low-energy irradiation, such as ultraviolet rays through grafting efficiency also increases with time for up to l a mercury lamp [28,29], high-energy irradiation, such h, then decreases. as y-rays from a CO-60source [30,31], and mechanical degradation by ball-milling of cellulose acetate in the 3. Effect of Hydrogen Peroxide Concentration presence of a monomer [32] are capable of generating Figure 5 shows the effect of hydrogen peroxide concenfree radicals on the cellulose molecules. Graftingof celtration. The grafting percentage and grafting efficiency lulosic material with monomer can be achieved by the as well as polymer loadingwere increased by increasing free radicalsmechanism,which can be produced by the dose of hydrogen peroxide added up to 0.5 mL/g physical or chemical treatment. y-irradiation is considcellulose. The same result was obtained for grafting of ered one of these treatments. The effect of ionizing rawoodpulpapplying the same method [27]. Upon indiation on the cellulosic fibers' properties has attracted creasing the hydrogen peroxide concentration over 0.5 considerable interest. It is generallyaccepted that oxidamL/g cellulose, the grafting percentage, grafting effition and cleavage are the principal reactions that celluciency, andpolymerizationefficiency are decreased. lose chains undergo on the exposure to ionizing radiation The initial increase in the graft yield is due to the in[33]. The possibility of modifying cellulosicfibers' propcrease in the number of xantb.ate groups that have reerties by graft copolymerization using this radiation has acted with hydrogen peroxide, thus creating more sites received considerable attention. Graft copolymerization (free radicals) that can initiate polymerization. is based on the existence of free radicals in an existing cellulose chains. The free radicals can be induced by 4 . Effect of Grafting Temperature various initiationmethods, such as ultraviolet and ionizing radiation [34]. Grafting of irradiated cellulose has The effect of the reaction temperature on the grafting process isdiscussed. Upon increasingthe reaction temlow-grafting efficiency, i.e., highquantity of homopolyperature from 15°C to 40°C the grafting decreases from mer. Figure 6 shows the effect of y-radiation dose on 30.7% to 29.1% and from 30.0% to 28.5% for cotton the graft yield percent of the grafted cotton linters with linters and viscose woodpulp, respectively. This can be MMA monomer. It is clear that the graft percent has a

89

83

116 1/3 1/2 2/3 1/6 m1 H 0 / g c e l l u l . 2 2

b

m1 H202

1/3 1/2 / cellul.

2/3

Figure 5 Effect of hydrogen peroxide concentration on grafting yield and grafting efficiency. 0 = cotton linters; 0 = wood pulp.

Nada and Yousef

534

10

8

4

0

10

40 60

30

20

dose

50

rad 1

( M

Figure 6 Effect of radiation dose on the graft yield (%) of grafted cotton linters.

KMn04 and different parameters affecting the grafting process, e.g., reaction temperature, reaction time, and monomer concentration, were investigated to ascertain the optimum conditions of the grafting process [23].

linear relationship with the increase of the y-radiation dose from 5-52 Mrad.

D. Grafting of Cellulose by Using KBrOa or KMn04 Redox System

1. Amount of Consumed KMnO4 Figure 7 shows the effect of KMn04 concentration on the consumed KMn04. The amount of KMn04 consumed by the cotton linters increases with increasing KMn04normality, i.e., MnOz deposit content enhances with increasing KMn04 concentration. In the presence of acid, free radicals are formed as a result of acid action on the deposited MnOz. These created free radicals depend onthe nature and normalityof the acid andproduce

Paper wood pulp was grafted with methylmethacrylate using the potassium bromate redox system. The effects of monomer concentration, reaction time, initiator concentration, and liquor ratio on the graft yield and polymer loading of the grafted samples were studied. The optimum conditionsof the grafting reaction were determined [24]. Cotton linters weregraftedwithmethylmethactylate usingthe KMn04 redox system. The effect of type and concentration of acid as the activator for

.

0 054 0 046

.

0 038

0.039

0.022

0.014 0

I

1

1

0.02

t

1

0.04 0.06 0.08 Normality of -no4

I

0.10

Figure 7 Effect of KMn04concentration on the amount of consumed KMn04. Liquor ratio1 :50; t = 30 min; T = 50°C.


535

cellulose macroradicals as follows: Cell-OH

+ R*+ C e l W + RH

(1)

In presence of the monomer cellulose, macroradicals are added to the double bond of the monomer.

2. Effect of Acid Concentration The effect of acid concentration on the polymer loading and grafting efficiency isillustrated in Fig. 8. The polymer loading increases with increasing acid concentration. The efficiency of MMA grafting onto cotton linters increases until the acid concentration of 0.05 N is reached. By increasing the normality of the acid more than 0.05 N, the grafting efficiencydecreases. This can be interpreted by the fact that higher acid normalityestablishes a balance between the suppression of the rate of formation of active species and the rate of generation of primary radicals. On the other hand, the decrement in grafting observed with oxalic acid is higher than that observed with sulfuric acid.This can bedue to thefact that oxalic acid is moreactive than sulfuric acidfor the formation of free radicals and, consequently, the grafting in the presence of oxalic acid isgreater than in the presence of sulfuric acid. W. ANIONIC GRAFTING Strongly acidic monomers, such as nitroethylene, can only be grafted under anionic conditions. Also, anionic grafting is useful for conventional monomers such as acrylonitrile andmethylmethacrylate. In additionto the

fact that certain monomers cannot be grafted by free radical methods, ioic grafting has some advantages, e.g., it is fast, it can be carried out at low temperatures, and homopolymer formation excluded. is Using this method, cellulose is treated with a Lewis acid, such as boron trifluoride, aluminiumchloride, or stannic chloride [35]. Initiation of graft copolymerization of vinyl monomers onto cellulose by formation of free radicals using ceric ions is the most widely used method. V. ACTIVATION OF CELLULOSE TOWARD GRAFTING

Since the cellulose hydroxyl groups may not be available for reaction, because of crystallinity or insolubility of the cellulose, this hinders the access of the reagent to the innermost hydroxyl groups. Nonuniform products can resultdue to the differences in accessibility of different portions of the cellulose. It is important, therefore, to have the cellulose in such a form that allows a maximum number of hydroxyl groups availablefor reaction. This can be broughtabout by decreasing the hydrogen bondingbetween the chains using various activation processes. Cellulose may be activated by various mechanical or physical treatments such as, beating, grinding, and swelling in different swellingagents. Different chemical treatments, involving the introduction ofnew groups into the cellulose chains, e.g., acetylation or carboxymethylation, may also be used for activating cellulose as a result of the opened crystalline structure of the lowsubstituted derivatives.

-

I

62

34

130

*

-26

4

P

22

58

- 18 -14

54

c

m

2

E

2 &

0 0.02

0.04

0.06

0.08

0.10

Normality of acid Figure 8 Effect of acid concentration on the grafting of cotton linters. [KMn041 = 0.06 N; [monomer] = 1 moYg cell; liquor ratio 1 :50; T = 50°C; t = 1 h. (-) Oxalic acid; (----) sulfuric acid.

Nada and Yousef

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A.

Mechanical or PhysicalActivation

1 . Beating The principal effect of beating is a physical one that causes breakage of some intrafiber hydrogen bonds and their replacement of fiber by water hydrogen bonds [36]. In addition, the graft yield and the number of grafted chains are greatly affected by the beating degree of the pulp [37], The effect of beating wood and bagasse pulps at various beating degrees on their reactivity toward grafting has been studied [38]. The beating process has a great effect on the properties of cellulosic fibers. It increases the fiber flexibility and increases the surface area byfibrillating the fibers and, consequently, increases the area of contact between the fibers. Beating causes some reduction in the size of the crystalline zones. Also, it can benoticed that minorchemical changes that occur in the beaten pulp may bedue to the large increase in surface activity and the changes in some other physical properties of the fibers. By comparing the effect of the degree of beating on the graft yield of pulp using a 50: 1 liquor ratio, 3 mL MMA: 1 cellulose, 1% acid for 2 h at room temperature, it isapparent that increasing the degree of beating up to 37"SR is accompanied by a remarkable increase in the 39% to 78%. This indicates that the reacgraft yield from tivity of the pulp toward grafting increases as the beating degree increases until reaching this optimum value. The reactivity tends to decrease when increasingthe beating degree more than 37"SR. The observed increase in the pulp reactivity caused by low-beating degrees is due to the significant increase in the surface area of the fibers and improved fiber flexibility. Beaten pulp of 61"SR shows the lowest grafting yieldat all monomerconcentrations. Total monomer conversion is increased by increasing the monomer concentration, and it decreases according to the following order: beaten sample for 37"SR > 20"SR > 48"SR > sample unbeaten (2)> 61"SR

2. Grinding In the case of grinding, the cellulose fibers go over a state of fine fibrillation into a more or less powdery substance. This mechanicalseverance of cellulose may break main valence bonds and will, therefore, decrease its degree of polymerization. In addition, the crystal structure of cellulose fibers is nearly lost 1321. Grinding of the cellulose fibers also, appreciably increases its surface area. Effect of Grinding on the Graftability Compared withuntreated pulp, a relatively high liquor ratio is necessary to attain maximumgraftyield (43.5%) because of the small particle size of the ground pulp. For this reason, ground pulpneeds a higher liquor ratio than the untreated pulp to help the mobility of the particles. On the other hand, comparing the graft yield

for ground untreated pulps at a liquor ratio 50: 1, it is clear that the graft yield attained by the ground pulp is higher thanthat of untreated pulp. This can beattributed to the high surface area of the ground pulp, which allows better contact between the pulp and chemicals.

3. Swelling The accessibility of cellulose for reaction is improved by usinga swelling agent, which expands the crystalline lattice and permits reagents to penetrate to all cellulose molecules. The action of the swelling agents in disrupting the intermolecular hydrogen bonds may be either reversible or irreversible. Water, amines, and aminesolutions have a reversible effect [40-441. Bleached cotton stalk pulp is treated with different concentrations of ethylene diamine (50-100%) for 20 min. It is clear that the crystallinity index (Crl) of these treated pulps is decreased by increasing the concentration of ethylene diamine; that is, the decrystallization increases. The degree of polymerization is nearly the same, but some increase is shown in the sample treated with 100% ethylene diamine. This indicates that 100% ethylene diamine may act as a dissolving agent for low degree of polymerization (DP) of cellulosic chains and hemicellulose. Of samples swollen withethylene diamine, the graft yield at a 50: 1liquor ratio increases as theconcentration of ethylene diamine increases. This is due to the increase of decrystallization of swollen samples, which helps the penetration velocity of the chemicals through the cellulosic chains. Graftability of the samples treated with 100% ethylene diamine is lower that of the sample treated with 75%. This is due to the dissolution of low DP chains and some ofthe hemicelluloses, which is detectable by the increase in DP of the sample teated with 100% ethylene diamine. From the preceding results, it is seen that there is a large differencein graft yield between grafted swollen pulp treated with ethylene diamine and untreated pulp at a liquor ratio of 50: 1 at all monomer concentrations. Summing up these results, it is clear that beating and swelling ofcotton stalk pulp have important effects ongraftability.Athigh-monomer concentrations, the graftability is higher in the swollen samples than in the beaten samples. B. Chemical Activation

It has been found that substitution of cellulosehydroxyls of cotton by acetyl groups, greatly affects the graft yield W]. The grafting reaction depends upon the degree of substitutionas well as the kind of pulp used.Introducing acetyl groups in the cellulose chains (high substitution) causes a large reduction of its swellability, which reduces the diffusion of the reactants. Thus, acetylation lowers the graftability of the cellulose.


537

Partial carboxymethylation of woodpulpsignificantly increases its susceptibility toward grafting with acrylonitrile usingthe ceric ion as the initiator [46]. Studies dealing with grafting of various vinyl monomers, such as acrylonitrile, methylmethacrylate, and acrylamide, onto partially carboxymethylated cotton cellulose using tetravalent cerium as the initiator have been reported [47]. Cotton linters and viscose grade wood pulp were partially acetylated and carboxymethylated. The Samples were grafted under the same conditions with acrylamide. Grafting yield and efficiency depend on several factors, such as the kind of pulp, chemical and physical structure, type of the introduced substituent, and degree of substitution (DS).

l . Grafting of Partially Acetylated Cellulose Wood pulp and cotton linters were partially acetylated to different acetyl contents ranging from approximately 6%-13% (0.23-0.54 degree of substitution [D.S.]). The increase of the DS of partially actylated cotton linters from 0.26 to 0.56 causes the percent graft of partially acetylated cotton linters to decrease slightly from0.93% to 0.6%. In the case of woodpulp, the effect is somewhat different-the percent graft increases with an increasing DS until about 0.48, then the percent graft decreases. Grafting of the unacetylated samples gives higher values in the case of wood pulp than in cotton linters. This is explained by the crystallinity of both pulps, which is 80% and 56.8% for cotton linters and wood pulp, respectively.

and the opening of the cellulose structure due to decrystallization. At an early stage of carboxymethylation of cotton linters, some of the OH groups are blocked, whereas the opening of the cellulose structure is very low or negligible, therefore, the graft yield tends to decrease. By increasing the D.S., the effect of the carboxyl groups in opening the cellulose structure begins to predominate and, hence, the graft yieldincreases gradually.

VI.

EFFECT OF GRAFTING MEDIUM

Graftcopolymerization of methylmethacrylate onto paper wood pulp using ceric ammonium nitrate as the initiator has beenstudied. Different experimental conditions have been used, including both water and waterorganic solvent systems. The effects of composition of the graftingmediumandgrafting temperature on the grafting process are examined. A.

Effect of Different Wetting Solvent-Water Systems

Table 1 shows that alcohols, e.g., methanol, ethanol propanol, butanol, andisobutanol, have a distinct effect of chain length on the grafting efficiency. There is a decrease in the graftingyieldwithanincreasingchain length and degree of branching of the alcohols. Copolymerization isthe most pronounced reaction in methanol-waterandis decreased in isobutanol-watermedium. A small molecule, such as methanol, is not only capable of swelling the trunk polymer but also miscible in all proportions with methylmethacrylate and water, thus simplifying access anddiffusion to the grafting sites. By contrast, butanol and isobutanolare relatively poor swellingagents for cellulose andthe grafting iscorrespondingly low (Table 1). By further increasing the methanol in the grafting medium, the graft yield decreases. This can be related to the lower solubilityof the initiator in the grafting medium and a reduced formation of free radicals, which

2. Grafting of Partially Carboxymethylated Cellulose In the case of cotton linters, the percent graft slightly decreases withincreasing D.S., then itbegins to increase. This behavior can be explainedby the opposite effects of introducing carboxyl groups. These effects are the blocking of some of the cellulose hydroxyl groups

Table l Effect of Ratio of Different Solvents in Water System on Polymer Loading and Grafting Yield

12/18

Solventlwater 8/22 4/26 ratio

0130 2/28

p.1. Typeofsolvent 32

54 Methanol 42 52 40 34 Ethanol 45 34 Propanol 43 32 20 Butanol 22 Isobutanol Water

1129

g.y. %

p.l. %

-49 34 46 32 3628 -25 25 22 21 43 33

40 42 35 30 28 25

%

p.1. = polymer loading; g.y. = grafting yield.

g.y. %

p.l. %

g.y. %

p.1. %

g.y. %

p.1. %

g.y. %

p.l. %

44 27

34

g.y. %

11

30

10

18

15

35 18 14

16 10 810

13 8 9

8 8

7 7 6

g.y. %

p.1. %

-

.

8 7

-

Nada and Yousef

538

decrease the initiation and propagation of the grafting process. The grafting in aqueous medium is higher than in case of the 97% methanol medium. This can be attributed to the lower dissolutionof the initiator in the solvent medium. Onthe other hand, graftingin methanol-water medium (8: 22) produces higher grafting than in water medium.

B. Effect of Mixed Solvent-Water Systems on Grafting

A mixture of methanol and butanol or isobutanol in water with a ratio of 4 :4 :22 in the grafting process produces a higher graft yield than a mixture of butanol or isobutanol in water with a ratio of 8 :22. Partial replacement of ethanol bymethanol has nearly no effect. In the case of propanol an increase in grafting is visible. This can be attributed to the mixing of higher carbon alcohols, e.g., butanol andisobutanol, with the active solvent methanol, which increases the miscibility of the monomer in these grafting systems and, consequently, increases the penetration of monomer to the active sites on the cellulose chains. C.

Effect of the Addition of Acetone on the Grafting Process

The effect of the addition of acetone to the grafting medium has been investigated.The acetone-water system produces a high-grafting yield, especiallyat low concentrations of acetone in the grafting medium. This is explained by the inhibition of the formation of homopolymer. By increasing the ratio of acetone-water up to 8 :22, the grafting yield is lower than that in case of the 8 :22 methanol-water medium. D.

Grafting in a Nonwetting Solvent

Grafting of methylmethacrylate onto cellulose using ceric ammonium nitrate (Ce4') as the initiator in a benzene-water system is also demonstrated. The grafting yield in a benzene-water system is much lower than in the case of the methanol-water system and decreases by increasing the ratio of benzene to water. This can be due to the lower polarity and wetting power of benzene, which leads to poor swelling of the cellulose.

ACKNOWLEDGMENTS

The authors wish to express their gratitude to Dr. A. Ibrahem for his kindcooperation and his helpful discussion of some parts of this work. The authors gratefully acknowledge the financial assistance received from cellulose of paper Det., National Research Centre, Dokki; Cairo, Egypt.

REFERENCES 1. A. Hebeish, and P. C. Mehta, Textile Res. J., 37 9115

(1967). 2. A. Hebeish, and P. C. Mehta, J . Appl. Polym. Sci., 12: 1625 (1%8). 3. H. A. Krassto, and V. T. Stannett, Adv. Polym. Sci.,4: 111(1965). 4. J. J. Guthrie, and Z. Haa, Polymer, 15: 133 (1974). 5. R. B., Phillips, J. Ouere, Z.Quiroy, and V. T. Stannett, Tappi 55: 858 (1972). 6. G. Mino, and S. Kaizerman, J . Polym. Sci.31: 242 (1958). 7. J. H. Richard, J . Soc. Dyers Colocirists, 80: 640 (1965). 8. R. W. Faessinger, andJ. S. Conta, U.S. Patent 3,359,224 (December 19, 1%7). 9. W.J. Briekman, Tappi, 56: 97 (1973). 10. G.Machell, and G. N. Richards, J . Chem. Soc., 3308 (1961). 11. S. Kaizerman, G. Mino and C. F.Meinheld, Textil. Res. J., 32: 136 (1962). 12. E. Schwob, V. T. Stannett, W. H. Rakowiz, and J. K. Magrane, Tappi, 45: 390 (1962). 13. R. M. Livshits, V. P. Alachev, M. V. Prokofeva, and Z. A. Pogovin, Vysekomolekul. Soedin, 6: 655 (1964). 14. S. Kimura, and H. Imoto. Makromole Chem. 42,140 (1960). 15. Y. Iwakura, T. Kutopaki, and Y . Imai, J . Polym. Sci., ( A ) 3: 1185 (1%5). 16. A. M. A. Nada, M. A. Yousef, and A. A. Ibrahem, Acta Polymercca, 38 93 (1987). 17. A.A. Shabaka, M. A. Yousef,and A. M. A. Nada, Polym. Plast. Techn. Eug., 29(1, 2): 167(1990). 18. A.M.A. Nada, S. F. El-Kalyoubi, and I. El-Roweing, PoIym. Plasr. Techn. Eug.. 2814): 439 (1989). 19. Q.Y. Mauson, Nagaty,A. Beshay, M. Nosier, J . Polym. Sci., 21: 715 (1983). 20. A. Hebeish, N. Abore-Zeid, A. Waly, and E. El-Rofie, J . Appl. Polym. Sci., 23: 306 (1979). 21. Y . Londier,and P. Label,GermanPatent 1,100,286 (Appl. Oct. 11, 1958). 22. A. A. Ibrahem and A.M. A. Nada, Acta Polymercca, 36: 342 (1985). 23. A. M. A. Nada, andM.A. Yousef, Acta Polymerica,40: 69 (1989). 24. M.A. Yousef, and A.M. A. Nada, Acta Polyanerica, 41: 581 (1990). 25. A. A. Ibrahem, and A. M. A. Nada, Bull., NRC, Egypt, 10: 206 (1985). 26. A. M. A. Nada, and A. A. Ibrahem, Acta Polymercca, 38 28 (1989). 21. S. El-Meadawy, and A. El-Ashmawy, Acta Polymerica, 34: 229 (1983). 28. I. Sakurad, T. Gkada, and K. J. Kaji, Polymer Sci., C37 1(1972). 29. J. J. Harris, J. H. Garra, I. V. Dogray, and J.C. Arthur, Jr. Text. Res. J., 42: 14 (1994). 30. T. Akira, Y. Sugahara, and Y. Honkawa, Transaction, 43(7): 362 (1987). 31. A. A. Shabaka, A. El-Agramy, and A. M. A. Nada, Isotopenpraxis, 27: 251 (1991). 32. W. Deters, and D. Huang, Fasuforsch. Textiltech, 19(S): 183 (1963). 33. J. G. Guthrie, M.B. Huang, and G. 0. Phillips, Eur. Polymer J., 8 47 (1982). 34. K. D. Lawrence, and D. Rades. Appl. Polym. Sci., 17 2653 (1973). 35. G. Rauseng, and S. Sumner, Tappi, 45(1): 203 (1%2).

Grafting of Cellulose 36. H. G. Higgurs, and J. de Yong, Transaction of Oxford SvmDosium. 2: 51 (1962). 37. Y.Ogrwara; and H:Kubota, Kogyo KagakuZasshi71(1): 171(1968). 38. S. Heileol, and S. El-Kalyoubi, J . Appl. Polym. Sci., 27: 3691 (1982). 39. S. Lipatou, D. V. Zharhovskii,and M. Zagraevskaya, Kollod Zh. 21: (1959). Absk Bull. Inst. Paper. Chem. 32(3), 1333 (1961). 40. N. I. Klenkova, Zur. Priki-Chim, 40: 2191(1967). 41.A. Koura, and A.A. Ibrahem, Faserfersch. Textiltech., 25: 57 (1974). 42.A. M. A. Nada, A.A. Shabaka, M. A. Yousef, and K. N. Abd El-Nour, J . Appl. Polym. Sci.,40: 731 (1990).

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35 Graft Copolymerization of Vinyl Monomers Onto Macromolecules Having Active Pendant Group via Ceric Ion Redox or Photo-Induced Charge-Transfer Initiation Kun-Yuan Qiu and Xin-De Feng Peking University, Beijing, China

1.

INTRODUCTION

initiation systems and their initiation mechanism; the synthesis of macromoleculeshaving active pendant group and its grafting reaction; and the grafting mechanism as well as the grafting site are reviewed in this chapter.

Graft copolymershave been synthesized from the grafting vinyl monomer copolymerization onto polymers by radical, anionic, and cationic initiation processing. In general, there are three ways for synthesis of graft copolymers by radical initiationreaction: (l) chain-transfer reaction; (2) redox initiation withceric salt, Ce(1V) ion, and photo-induced charge-transfer reaction; and (3) copolymerization of vinyl monomer with macromer[l ,2]. However, these methods yield a mixture of homopolymer, graft copolymer, and ungrafted original polymer. Some systems such as high-impact polystyrene (HIPS), acrylonitrile-butadiene-styrene (ABS) polymer, and methyl methacrylate-butadiene-styrene (MBS)polymers have been commercialized. For characterization, the purification of graft copolymer necessary is and even the purification process is tedious. Furthermore, the free radical graftingreactions of polymers can proceed either along the main chain(backbone) of the polymer or from the pendant groups (side chain), these make it difficult to realize the grafting site in a graft copolymer. In our laboratory on the basis of investigation of vinyl radical polymerization initiationsystems [3-71, i.e., Ce(1V) ion redox and amine-benzophenone photo-induced chargetransfer initiation systems, wehavedevelopedsome techniques for graft copolymerization via initiation reactionof active pendant groups of polymers either by redox or photo-induced charge-transferreaction. The results of our research work concerning the model compounds of active functional group for Ce(1V) ionredox

II. CERIC ION REDOX INITIATION SYSTEMS

Since the 1950s, ceric salt, Ce(1V) ion alone, or Ce(1V) ion redox systems consisting of Ce(1V) ion andalcohol, aldehyde, ketone, amine, and carboxylic acid were investigated to initiatevinylmonomers to polymerize [S-121. Although somepapers reported the grafting reaction of polymer initiated by Ce(1V) ion [13-161, much less attention was paidto thegrafting reaction of macromolecules having active pendant groups initiated with Ce(1V) ion. Lin et al. [17,18] have reported the graft copolymerization of acrylamide ( A A M )onto macromolecules having active pendant groups such as 3,4-dihydroxycyclohexylethyl acrylate and2,3-dihydroxypropenyl acrylate initiatedwith Ce(1V) ion.In order to develop the scope of the active pendant group, a series of works concerningthe effect of model compounds on the rate of polymerization in the presence of Ce(1V) ion were carried out in our laboratory, and some novel redox initiation systems consisting of Ce(1V) ion-acetanilide [3,4], Ce(1V) ion-alkyl tolylcarbamate [19], Ce(1V) ion1,3-dicarbonyI[20,21],Ce(1V) ion-triketone [22], Ce(1V) ion-actoacetanilide [23], and Ce(1V)ion-N-acetyl-N'-4541

Qiu and Feng

542

tolylurea [24] have been established. The initiation mechanism wasproposed based onthe end group analysis of the resulting polymer andelectron spin resonance (ESR) studies. Moreover, the active functional group mentioned previouslycan serve as an active site for the grafting reaction of polymers havingsuch an active pendant group. Grafting reactions of macromolecules having 4-tolylcarbamoyl, 1,3-diketone, and 4-tolylurea pendant groups initiated with Ce(1V) ion were performed successfully and are discussed in Section V. 111.

NOVEL CERIC ION REDOX INITIATION SYSTEMS

Recently, in our laboratory the following novel Ce(1V) ion redox initiation systems have been investigated for vinyl radical polymerization. A.

Ce(lv)Ion-Phenylcarbamoyl Compound Systems

Samal et al. [25] reported that Ce(1V) ion coupled with an amide, such as thioacetamide, succinamide, acetamide, and formamide, could initiate acrylonitrile (AN) polymerization inaqueous solution. Fenget al. [3] for the first timethoroughlyinvestigated the structural effect of amide on AAM polymerization using Ce(1V) ion, ceric ammonium nitrate (CAN) as an initiator. They found that only acetanilide(AA) andformanilide (FA)promote the polymerizationandremarkably enhance R,. The others such as formamide,N.N-dimethylformamide (DMF), N-butylacetamide, and N-cyclohexylacetamide only slightly affectthe rate of polymerization. This can be shown bythe relative rate (Rr),i.e., the rate of AAM polymerization initiated with ceric ion-amide divided by the rate of polymerization initiated withceric ion alone. R, for CAN-anilide system is approximately2.5, and the others range from 1.04- 1.1 1.

I . Ce(IV) Ion-AcetanilideSystems Feng et al. [3] have studiedthe structural effect of acetanilide on the AAM polymerization either in water-formamide [31, water-acetonitrile 141, and water-DMF [26] mixed solution using Ce(1V) ion-acetanilide and its substituted derivatives as the initiator. The results showed that an electron donating substituent on the phenyl group would enhance the R,, while an electron withdrawing group would decrease it, as shown in Table 1

MI. 2 . Ce(IV) Zon-Alkyl Phenylcarbamate Systems We have investigated Ce(1V) ion-carbamates, such as methyl and butyl4-methylphenyl-carbamate, (MTC and BTC, respectively), or methyl, ethyl, and butyl phenylcarbamate, (MPC,EPCandBPC, respectively), systems for AAM polymerization [191. It was found that the presence of carbamate compounds can promote the polymerization andenhance the rate of AAM polymerization (R,) in descending order as: MTC = BTC MPC = EPC = BTC (1) The structural effect of alkyl groups such as methyl, ethyl, and n-butyl on the R, is small. Alkyl 4-methylphenylcarbamate can be chosen as a model compound for the hardsegment of poly(ether-urethane) (PEU). This group caninitiate grafting reaction with Ce(1V) ion and the grafting site was proposed at the hard segment of PEU [3,15] as shown in Scheme (1). - ---- ---- - ---.--,

*

c

0

I

0 :

MTC

Scheme 1

Table 1 Structural Effect of Acetanilide on AAM Polymerization

Compound

R, x lo5 Relative (moYL.s)

Substituent (X)

15.3 19.3

0.76

AA p-APT m-AABA m-AA1 m-AAe 0-AABA 0-AA1 o-AAe

H1.92 2.97 m-COOH 2.06 m-CH20H m-CH20COCH3 0-COOH 0-CH2OH 1.46 o-CHZOCOCH~

8.75 16.8 26 18 17.5 0.77

13 12.8

rate R,

R(m)

1.oo 1.oo 1S 4 2

1.04 1.14

1.48

AA = acetanilide; p-APT= N-(p-methylphenyl) acetamide; m-AABA= m-acetamino benzoic acid;m-AAI = m-acetamino phenylmethanol; m-AAe = o-acetaminophenylmethyl acetate; o-AABA = o-acetaminobenzoic acid; 0-AA1 = o-acetaminophenylmethanol;o-AAe = o-acetaminophenylmethyl acetate.

Ceric

Initiation Systems

543

B. Ce(lv) Ion-l ,&Dicarbonyl Compound Systems

possess higher activity than alcohol compounds, in other words, the redox reaction between aldehyde and Ce(1V) ion is morefavorable than alcohol. Therefore, based on the investigation of the effect of cis- and trans-l ,2 cyclohexane-diolon the AN polymerizationinitiatedwith Ce(1V) ion, Li et al. [30], and Yu et al. [31] commented that the following Mino's mechanism (1956) of 1,Zdiol with Ce(1V) ionas a fashion for the grafting reaction of cellulose and starch is questionable:

Although carbonyl compounds, such as formaldehyde [27,28], can couple with Ce(1V) ion to initiate acrylonitrile (AN) or methyl methacrylate (MMA) polymerization, the remarkable activity of aliphatic aldehyde had not been noticed until the paper of Sun et al. [29] was published. They found that aliphatic aldehydes always "CH-CH-

I

OH

I

+

CF

-

-CH

II

0

OH

+

.CH -

I

cr

+

+

H+

OH

IM

Polvmerization 7 Yu et al. [31] pointed out that hydroxymethene radical will react with Ce(1V) ion to form a carbonyl group very quickly, and then can it further oxidize to form carbonyl radical. Therefore, the grafting polymerization would take place on the carbonyl radical as the following:

O=CHCH

I

PM

II

0

o=c

I

PMPM

c=o I

Scheme 2

Voong et al. [32] have investigated the effect of ketone compounds on AN polymerization in the presence of Ce(1V) ion and found that acetone, 2-butanone, 3-pentanone, cyclopentanone, and cyclohexanonewere effective in promoting polymerization. Much less attention has been paidto the Ce(1V) ion-1,3-diketone compound initiation system. In a patent [33], an initiator consisting ofCe(1V) ion-acetylacetone(2,4-pentanedione, AcAc) was used for AAM polymerization. To date no detailed information has been published. Recently, we a series ceric [20,21,34-361 have thoroughly investigated redoxinitiation system containing1,3-diketonecompoundsandfound that 1,3-diketones, such as AcAc, benzoylacetone (BzAc), 3-benzyl acetylacetone(3-benzyl-2,4-pentanedione,[BzyAcAcl), dibenzoylmethane (DBzM), andtriketone, such as 3-benzoylacetyl acetone (BzAcAc), and 1,3-ketone ester, such as ethyl acetoacetate (EAcAc), coupled with Ce(1V) ion can aform redox initiator for AAM polymerization. The data are compiled in Table 2. It can be seen that all diketones, even the substituted diketone such as BzyAcAc, can remarkably enhance the R,. BzyAcAc can be chosen as a model of poly[3-(4-vinylphenylmethyl)-2,4-pentanedione] [P(St-

Table 2 Effect of 1,3-Diketones or Triketone on AAM Polymerization Initiated by Ce(IV) Ion R~ X 104

Reductant

None 2-Butanone AcAc 9.43 CH,COCH(CHzCsHs)COCH3 BzyAcAc ~COCH3 BzAc HzCOCsH5 DBzM 3COCH(COCaHs)COCH3 BzAcAc

(m0VL.s)

CH~COCHZCH~~ CH3COCHzCOCH3

[MM1 = 1.0 m o l L ; [CAN] = 2.5 X [MM1 = 2.0 m o a ; [CAN] = 5.0 X HzO/CHXN = 3: 1 ( v h ) at 40.0"C. CAN = ceric ammonium nitrate.

1.2 2.37

m o m ; [Reductant] = 1.0 x 10-3 morn. m o a ; [Monoketone] = 5.0 x 10-3 m o a .

E, (kJ/mol) 41.2

-

Qiu and Feng

544

CH*-AcAc)],the grafting reactions of this polymer initiated withCe(1V)ionwill be discussed in Sec. V1.B. Furthermore, the order of the promoting activity of diketone is always greater than that of monoketone.

C. Ce(lv) Ion-AcetoacetanilideSystems As wehave mentioned previously,1,3-diketoneand anilide were very effective reducing agents for vinyl polymerization initiated by ceric ion, respectively. Acetoacetanilide (AAA), a compound having a 1,3-diketone and an anilide structure as well: I

R :

anilide

CH3-C "CH2-C I

II

! I

- N H a I

Canit promote vinylpolymerizationinitiatedwith Ce(1V) ion? Donget al. [37-391 for the first time reported that AAA and its derivatives such as o-acetoacetotoluidine(AAT), o-acetoacetanisidide (AAN),and2-benzoyl acetanilide (BAA) possess very high reactivity toward Ce(1V) ion ininitiating the polymerization of vinyl monomer. The results are tabulated in Table 3. From Table 3, it can be seen that the reactivity of acyl acetanilide, such as BAA or A A A , is higher than that of the other reductant reported from our laboratory, N-acetyl-p-methylaniline ( p i.e., acetanilide(AA), AFT),acetylacetone (AcAc),and ethyl acetoacetate (EAcAc). Moreover, the promoting activities of derivatives of acetoacetanilidewere affected bythe ortho substituent in benzene ring, andthe relative rate of polymerization (R,) decreased with the increase of the bulky ortho substituent to the redox reaction between Ce(1V) ion and substituted acetoacetanilide.

D. Ce(lV) lon-N-Acyl-N'-6tolylurea Systems

Urea has no promoting effecton the AAM polymerization initiated with Ce(1V) ion. Recently, Qiu et al. [24] havestudied the effect of N-acryloyl-Nf-4-tolylurea (ATU), N-methacryloyl-N'-4-tolylurea(MTU), and N acetyl-Nf-4-tolylurea(AcTU) on M M polymerization initiated with Ce(1V) ion and found that these three urea compounds havea high promoting effect on the polymerization of AAM. The data are cited in Table 4. It can be seen from Table 4 that the descending order of RP is as follows: CAN-AcTU > CAN-ATUCAN-MTU

(2) It reveals that either the steric effect of the bulky N substituent on the interaction of Ce(1V) ion and urea reductant, or the electron withdrawing effectof the vinyl group, will reduce the coordination of Ce(1V) ion with the carbonyl group, thus resulting in a decrease of the RP. IV.

INITIATION MECHANISM

To clarify and reveal the initiation mechanism of the Ce(1V)ion redox system in vinyl polymerization, we have thoroughly determinedthe initial radical structure by means of ESRspectrum analysis and end group analysis of the resulting polymer. The ESR spectra can reveal the structure of the free radicals directly. Sometimes, however, it is difficult to trap a nitrogen center radical in an amide, such as CHsCON.(CaH5) with 2methyl-2-nitrosopropane (MNP), withthe result that no signal is observed in the ESR spectrum of CAN/AA/ MNP system. End group analysis isa unique technique to reveal the presence of reductant moiety as an end group in the resulting polymer initiated withthe Ce(1V) ion reductant system. This methodalso can demonstrate the structure of the initial radical indirectly-asthe radical formed, it will initiate monomer polymerization and will enter into the polymer chain as an end group.

Table 3 Effect of Acetoacetanilide and Its Derivatives on AAM Polymerization Initiated by Ce(IV) Ion in Aqueous Media at 25.0"C

R~ X 104 (molhs)

Reductant

-

-

AAA AA" AAN BAA [CAN] = 1

C6H5NHCOCH2COCHs 30.3 O"X3CaH4NHCOCH2COCH3 24.0 O-OCH3C6H4NHCOCH2COCHs 22.3 C6H5NHCOCH2COC6H5 34.9 X

m o a ; [Red.]

= 0.5 X

morn; [AAM] = 1.0 mom.

1.6 48.4 38.4 35.6 55.9

R, 1.o

E, (kJ/mol) 58.3 9.18

- 1.26 -

Ceric Ion Redox initiation Systems Table 4 Effect of N-acyl-N'-tolylurea on AAM Polymerization Initiated by Ce(1V) at 35°C in H2O-THF (v/v = 1:1)

RP x lo4 (moVL.s)

Structure Reductant

-

No polym. during 1h ATU CHdHCONHCONH-CaH4CH3-p 0.86 MTU CH~"C(CH~)CONHCONH-C~H~CH~-P 0.78 CH3CONHCONH-CsH4CH3-p AcTU 1.24 [CAN]

=

5.0 x

m o a ; [Red.] = 1.0 x

The initialradicalsformedfromthereactionof Ce(1V)ionand reductant systems can be trapped by MNP. The spin adducts of the initial radicals and MNP were observed by means ofan ESR spectrometer. The structure of the initial radicals the and hyperfine splitting constants of the spin adduct of the radical with MNP and a-phenyl-N-tert-butylnitrone(PBN) are compiled in Tables 5 and 6, respectively. The formation of spin adducts from radicals and MNP is as follows:

m o a ; [AAM] = 0.2 m o a .

The formation of spin adducts from radicals and PBN is as follows: R - + (CH3)3

-29-=CHCaHs

-

( C H ~ ) J C" N - C H C ~ H S

4 k

The initialradicalsformedfrom the Ce(1V)ion redox systemcan initiate a monomerto polymerize and form an end group of the resulting polymer. When the reductant exhibits a carbonyl group, the amide group can be conveniently detected by the FT-IR spectrum of the polymer, such as polyacrylonitrile (PAN). The FT-

Table 5 Hyperfine Splitting Constants of CAN-Reductant-MNP Systems [20,36,38,40]

Hyperfine splitting constant (mT) MNP bytrapped Reductant Radical

AcAc BzyAc BzAc DBzM H*(COCaH5)aBzAcAc COOC2H5) EAcAc BAA

H5)

545

AAA

AAT AAN

(CHsCOhC.(CH2CsHs) CH3COCH.(COC6H5) 0.46

0.63 0.48 0.36 C~HSNHCOCH.(COC~H~) 0.49

CaHsNHCOCH.(COCHs) o-CH3C6H4NHCOCH.(COCH3) 0.09o-OCH3CaH4NHCOCH.(COCH3) 0.22

0.09

&xN

asN

a,"

1.36 1.44 1.41 1.41 1.4 1.39 1.45 1.43 1.44 1.43

0.36

-

-

0.3 0.25

A triketone compound BzAcAc when it reacts with Ce(1V) ion, the free radical observed is obtained BzAcAc.

a

0.09

splitting an acetyl group in

Table 6 Hyperfine Splitting Constants of CAN-reductant-PBN Systems [38]

Hyperfine splitting constant (mT) PBN bytrapped Reductant Radical

&xN

BAA AAA

AAT AAN

CH3COCH2CON*(C6Hs) 0.3 CH~COCH~CON.(C~H~CH~-O)

CHJCOCH~CON*(C~H~OCH~-O) 0.23

1.44 1.43 1.44 1.43

asN

arN

0.49

0.07 0.09 0.09 0.09

0.25

546

Qiu and Feng

IR spectra of PAN obtained from the CAN-1,3-diketone systems revealed that, besides the characteristic absorption band of CN group at 2244 cm", the characteristic absorption bands at 1727 and 1700 cm" (keto form of 1,3-diketone) and1625 cm" (enol formof 1,3-diketone) were observed simultaneously.Therefore, the 1,3-diketone moiety is present as an end group in PAN. Although the initial radical formed from the reaction of C e O ion and acetylanilide (AA) and N-p-tolylacetamide (PTA) hasnever been observed in the ESR studies, the presence of AA, PTA moieties inthe end group of PAN obtained from initiating the CAN-AA, CANPTA system have been detected by the FT-IR spectra analysis method. Similar results were observed in the end group analysisof CAN-phenylcarbamate, CAN-Nacyl-l\"-tolylurea initiation systems. Based on the ESR studies and the end group analysis, the initiation mechanism of Ce(1V) ion redox systems is proposed as: Ce(1V) ion-acetanilide system: Cc? + H20 === tCe(0H)I)'

+

r

R

l 3+

H

3+

J

P

"

- .. ..".c

H

-

Ce(1V) ion-l ,3-diketone system: Ce4' + Hz0 r C ~ O H 1)"

Scheme 4

+

CC?

+

H20

= r c w w )13'

CdOH) 1" + RCOCH2CONHAr

Complex I

-

02 II R-C

+

i"" " Complex 2 - 3+

COH

, H 'c CH

-

LONHAr

Complex2

-

H+ Complex 1

RCOkHCONHAr

lM

Polymerization

c I

CHTC, &OR

hr

Polymerization

Scheme 5

V. GRAFTINGREACTION OF MACROMOLECULES

H* l

Ce(1V) ion-acetoacetanilide system:

Among the transition metal ions, Ce(1V) ion isthe most widely used initiator for grafting reaction of cellulose. Many papers reported on the grafting copolymerization of vinyl monomers onto cellulose by using the Ce(1V) ion as the initiator. This technique is based on the fact that when cellulose is oxidized by the Ce(1V) ion, free radicals capable of initiating vinyl polymerization are formed on the cellulose. Since the radicals on the cellulose are produced more quickly than the homopolymerization of monomer initiated by the Ce(1V) ion alone, the result is that less homopolymers are obtained. Some review articles [13,41,42] have summarizedthe graft copolymerization usingthe ceric ion method. Although the mechanism of grafting copolymerization is still not completely confirmed, it is important to note the work of Yu et al. [31] on the studies of grafting mechanism of 1,2diols compound, as pointed out in Section 1II.B.

H*

VI. GRAFTINGREACTION OF MACROMOLECULES HAVING AN ACTIVE PENDANT GROUP INITIATED WITH CERIC ION Lin et al. [17,18] have reported the grafting copolymerization of AAM onto macromolecules havingactive pendant groups using the Ce(1V) ion as the initiator. Few studies in the similar field of grafting reaction are reported so far. Based on the investigation of the novel Ce(1V) ionredox system, we have developed a new technique for the grafting reaction of macromolecules having an active pendant group, such as p-tolylcarbamoyl(4-

Ceric

547

methylphenylcarbamoyl)and1,3-diketone pendant group, initiated with Ce(1V) ion. The synthetic methods of macromolecules having an active pendant group include:(1) the transformation reactions of polymer and copolymers, and (2) polymerization and copolymerization of functional monomers having active pendant groups. The macromolecules, either in the shape of film or microbeads, can be used as the substrate. As we have mentioned previously, the rate of polymerizationinitiatedwith the Ce(1V)ion redox systemis much faster than that initiated by Ce(1V) ion alone, as expressed in R, B 1. Therefore, the graft copolymerization is favored much morethan the homopolymerization, and the yield of the homopolymer was depressed to a very low level.

A. Grafting Reaction of Macromolecules

As mentioned previously, the 4-tolylcarbamoyl group is highly reactive toward the Ce(1V) ion. The graft copolymerization of AAM onto the films of copolymers 1 6 in the presence of Ce(1V) ion would take place, and the order of grafting percentage (G%) wasobserved as:

of the functional monomer unit the viahydrogen abstraction of amide to form a nitrogen center radical capable of initiating graft copolymerization. Other types of copolymers having a 4-tolylcarbamoylpendant group, poly(VAc-co-MAMT) 7 andpoly(MA-co-MAMT) 8, were synthesized from the reacacetate-co-maleic tion of 4-toluidinewithpoly(viny1 anhydride)andpoly(methy1acrylate-co-maleicanhydride), respectively[44].Graftcopolymerizationof AAM onto films of copolymers 7 and 8 were successfully performed by using the Ce(1V) ion as the initiator. The grafting reaction created from the reaction of the tolylcarbamoyl group withthe Ce(1V) ion to form the nitrogen radicalactive site to initiate AAM graft copolymerization was proposed as follows (Scheme 6):

-

1>2%-3>4%5>6

(3)

Obviously, the copolymers containinga more reactive functional monomer, i.e., monomer having a 4-tolylcarbamoyl group, have a higher graftingpercentage. The grafting mechanism is proposed similar to the initiation mechanism of the Ce(1V) ion and acetanilide system, and the grafting reaction occurred at the pendant group

Having 4-Tolylcarbamoyl Pendant Group

Macromolecules having 4-tolycarbamoyl and phenylcarbamoyl pendant groups can be synthesized by the polymerization and copolymerizationof functional monomers. Qiu and Song [43] have reported that the copolymers having 4-tolylcarbamoyl and phenylcarbamoyl pendant groups were synthesized by the copolymerization of functional monomers, i.e., N-4-tolylrnethacrylamide (NTMAAM), N-4-tolyl-acrylamide (NTAAM), and N-phenyl-methacrylamide (NPMAAM) withmethyl acrylate (MA)andmethyl methacrylate (MMA), respectively. The copolymers are represented as follows:

Q (Scheme 6 continues)

Qiu and Feng

548

Scheme 6

Recently, poly(itac0namide) with 4-tolylcarbamoyl pendant groupshave been synthesizedin our laboratory. The polymer 9 and copolymers 10 and 11 were synthesized via aminolysisof poly(N-4-methyl-phenylitaconi-

mide) and its copolymers, respectively. Graft copolymerization of M M on the surface of the polymer films using Ce(1V) ion as the initiator has been investigated by us [45]. The mechanism of graft copolymerization was proposed as shown in Scheme 7.

Q

y f 0,R = -CH3, 10

U

Graft copolymerization Scheme 7

The mechanism proposed that the nitrogen radical was formed in the 4-tolylcarbamoyl groups rather than in the butylcarbamoyl groups, which were both pendant groups of the polymers. In our previous work, as mentioned in Section 111, the promoting .activity of the 4tolylcarbamoyl group (inAPT R, = 17) is much higher than the butylcarbamoyl group (inN-butylacetamide R, = 1.10). Therefore, we proposed that graft copolymerization took place predominantly on the 4-tolylcarbamoyl pendant group.

Although we have proposed the grafting reaction mechanism of lpolymers having 4-tolylcarbamoyl pendant group initiated with ceric ion and the grafting site of the copolymer at the N-H of the 4-tolylcarbamoyl linkage via the oxidation of H by Ce(1V) ion as shown in Scheme (3), there is still lack of evidence for the grafting site, due to the difliculties of analyzing the nitrogen radicalsignalsin the initiation system byESRmeasurement. Recently, Zhaoand Qiu[461 havedesignedtwo

Initiation

Ceric Ion Redox

549

kinds of polymershaving4-tolylcarbamoyl pendant poly(4-acetaminostyrene), P(%groups, i.e., NHCOCH3), poly(NTMAAM)12, poly(NTMAAM-coMA) 13, poly(NTMAAM-co-EA) 14, andpoly(NTMAAM-co-VAC)15, and they werechosen as the original polymer in the grafting reaction withCe(1V) ion. Poly(4-acetaminostyrene) was prepared from crossI

-CHrCHNaOH

linked polystyrene beads by nitration, reduction, and acetylation of the three-step reactions. Acrylamide graft copolymerization tookplace onto poly(4-acetaminostyrene) beads initiated with Ce(1V) ion, and the graft copolymer P[(St-NHCOCH&g-AAM] was characterized. The mechanism of graft copolymerization is shown as Scheme 8 .

-CHTCH-

Q

+ CHTC-ONa

bl

When P[(St-NHCOCH3)-g-AAM] was hydrolyzed in the basic solution noPAAM was released. The scanning electron microscopy (SEM) micrographof the copolymer shows that the hydrolyzed grafted beads are still covered with PAAMs with salient micrographs. The results reveal that AAM graft copolymerization is initiated by the nitrogen radicalrather than any other radical. In a comparable experiment, AAM graft copolymers wereprepared by the AAM graft copolymerization

on the surface of polymer filmsof polymers 13-16 using the Ce(1V) ion as the initiator. Hydrolysis of the grafted copolymer filmsin NaOH aqueous solution caused the grafted PAAM macromolecules to separate from the substrate films. The hydrolysis of the grafted copolymer provides some evidence for the grafting site and the mechanism of graft copolymerization of polymers having a 4-tolylcarbamoyl group. The reactions are shown in Scheme (9).

NaOH

c=o bNa

Scheme 9

B. Grafting Reaction of Macromolecules Having 1,3-Diketone a Pendant Group

molecules having a 1,3-diketone pendant group such as poly[3-(4-vinylphenylmethyl)-2,4-pentanedione], P(StCH2-AcAc) 16 and poly[N-(4-acetoacetylphenyl)-methRecently, Zhao et al. [34,471, andQiu et al. [211 haveacrylamide], systematic name:poly(l-[4-(2-methyl-l0x0 - 2 - propenylamino]phenyl- 1,3 - butadiene} pMPApB) reported the graft polymerization of AAM onto macro-

Qiu and Feng

550

17, were initiated by the Ce(1V) ion, respectively. The formation of the grafted copolymer was revealed by x-ray photoelectron spectrometry (XPS) spectra, FT-IR spectra, and photomicrographs. Based on the ESR studies of Ce(1V) ion-BzyAcAcMNP, Ce(1V) ion BzAc-MNPsystems as mentioned before, the grafting reaction of P(St-CH*-AcAc)will take place on the methene carbon of 1,3-dikeone viathe abstraction of hydrogen by the Ce(1V) ionto form radicals and then initiate monomer graft copolymerization. The initiation mechanism of graft copolymerization is proposed in Scheme (IO).

methene carbon of 1,3-diketone by the abstraction of hydrogen with Ce(1V) ion as shown in Scheme (11). CH:, I

c=.

4H2-C-

CY I

Cc"

4 H r C -

+

C2

+

H'

Lo

Scheme 11 16

C. Grafting Reaction of Macromolecules

Graft copolymerization

Having 4-Tolylureido Pendant Groups

Scheme 10

In P(MPAPB) there are two types of functional groups, i.e., 1,3-diketone group andmethacrylanilide group, whichcan probably react with the Ce(1V) ion to formradicals to initiate the graftcopolymerization. From the kinetic studies, we have found that 1,3-diketone group hasa much higheractivity than that of acetanilide group containing electron withdrawing substituent acetyl at dposition of the benzene ring. Therefore, the grafting reaction of P(MPAPB) willtake place onthe

Qiu et al. [24]have reported the synthesis of macromoleculeshaving Ctolylureido pendant groups, such as poly(N-acryloyl-N'-4-tolylurea-co-ethy1acrylate) [poly(ATU-co-EA)]18, and poly(N-methacryloyl-N'-4-tolylurea-co-EA) [poly(MTU-co-EA)]19, from the copolymerization ofATUandMTU with EA, respectively. Graft copolymerization of acrylamide onto the surface of these two copolymer films took place the using Ce(1V) ion as initiator. The graft copolymerization is proposed as Scheme (12). "CH

rFHl 0

."/-CHrC-

k * H 5

R = H.

18

CH3.19

R I

L I

+

Cc* + H+

NH L O

.A

k

1Gran copolymcrizalion

Scheme 12

l . Characterization of Graft Copolymers

the higher water absorbability of PAAM. XPS determination of the surface composition of polymer films When the AAM monomer was grafted onto polymerwouldgive the higher ratio value of NIC on the surface films, the grafted films showed much higher water abof grafted films than that on the surface of the original sorption than the original ungrafted films. This is due to films, SEM photographs of grafted films always showed

551

Ceric Ion Redox Initiation Systems

salient micrographs, showing the presence of the PAAM macromolecules chain on the surface of the films, while the originalfilms were smoothand essentially featureless.

2. Advantage of Grafl Copolymerization Initiated with Ce(IV) Ion As we have previouslymentioned, the graft copolymerization of vinyl monomer onto macromolecules having an active pendant group initiated with the Ce(1V) ion will give the following features: (1) Since the radicals capable of initiating vinyl graft copolymerization are producedalmostexclusivelyon the active pendant group, grafting is favored much morethan homopolymer formation as compared withother grafting reaction systems; (2) The redox reaction of an active pendant group with Ce(1V) ion is muchfaster than the Ce(1V) ion initiating vinyl homopolymerization, and graft copolymerization will usuallytake place rather than homopolymerization; (3) The homopolymer yields can be depressed to very low levels and the separation of homopolymer from the grafted films is very easy. D. Mechanism of GraftCopolymerization of Chitosan Model Compounds

Chitosan, having a similar chemicalbackbone as cellulose, is a linear polymer composed of a partially deacetylated material of chitin [(1-4)-2-acetamide-2-deoxy-fl-~glucan]. Grafting copolymer chains onto chitosan can improve some properties of the resulting copolymers [48-501. Yang et al. [l61 reported the grafting reaction of chitosan using the Ce(1V) ion as an initiator, but no detailed mechanismof this initiationhas been published so far. Recently, Li et al. [30],Yu et al. [31]reinvestigated the mechanism of graft copolymerizationof vinyl monomers onto carbohydrates such as starch and cellulose initiated by the Ce(1V) ion with some new results as mentioned in Section 11. Furthermore, they investigated the mechanism of model graft copolymerization of vinyl monomers onto chitosan [51]. They chose the compounds containing adjacent hydroxyl-amine structures, such as D-glucosamine, trans-2-amino-cyclohexanol,2CH7

amino-3-phenyl-l-propanol,l-amino-Zpropanol, as the model for chitosan reacted with the Ce(1V) ion. Based on experimentalresults, they suggest that there are two ways for the initiation reaction to take place, depending on the reaction temperature. Atlower temperatures (40°C), the adjacent hydroxyl-amine linkage was oxidized to aldehyde and -CH=NH groups. At higher temperatures (90”C), however, the -CH=NH group could be hydrolizedto form another aldehyde group. In each case the aldehyde group was further oxidized to an acyl radical, therefore, one or two acyl radicals were responsible for initiatingpolymerization at 40°C and W”C, respectively. From this initiation mechanism,the important role of the aldehyde group in the reaction mechanism of 1,2diol and l-amino-2-hydroxy compound and Ce(1V) ion initiation systems can again be seen. VII.GRAFTINGREACTION OF MACROMOLECULES HAVING AN ACTIVE PENDANT GROUP BY PHOTOINDUCED CHARGE-TRANSFER INITIATION

A convenient method for synthesizing macromolecules that have amino pendant groups is by means of radical polymerization and copolymerization of the functional monomer containing an amino group, the so-called polymerizable amine. A well-known photo-induced charge-transfer initiation system is the benzophenone (BP)-amine system, which can be used as the initiator for photopolymerization of vinyl monomers[52-581. Li et al. [521 investigated the macromolecules having an amino pendantgroup, i.e., copolymer of N-(4-dimethylaminophenyl) acrylamide andMMA[poly(DMAPAACO-MMA)201. In the presence of BP under irradiation of ultraviolet light, the active amino group would react with BPto generate two kindsof radicals, i.e., the anilomethyl radical, which wouldinitiate MMA for graft copolymerization,and the semipinacol radical, which would undergo primary termination. Therefore, the graft copolymer was obtained without contamination of homopolymer PMMA (Scheme 13).

Qiu and

552

Similar results were obtained in the photografting reaction of N-methyloyloxethyl-N-methyl-p-toluidine

polymer[(PMEMT)211 and its copolymer, poly(MEMTco-MA) 22, using BP as the initiator [53], as shown in Scheme (14).

IM

Graft copolymerization

Scheme 14

Recently, Si et al. 159,601have investigated the synthesis of polymerizable amines, such as A"(3-dimethylaminopropyl) acrylamide(DMAPAA) and N-(3-dimethylaminopropyl) methacrylamide (DMAPMA), and their copolymerization reaction. DMAPAA or DMAPMA in conjunction with ammonium persulfate was used as a redoxinitiator for vinylpolymerization.Copolymers havingamino pendant groups, such as copolymer of

MMA andDMAPMApoly(MMA-co-DMAPMA) 23, obtained by radical copolymerization, can produce a photografting reaction with acrylonitrile (AN) usingBP as the initiator [61]. The formation of a graft copolymer, poly[(MMA-co-DMAPMA)-g-AN] wasconfirmed by FT-IR spectrophotometry. Based on ESR studies and end group analysis, the mechanism of grafting reaction is proposed as follows:

I"

Gran copolymerization

Scheme 15

Therefore, the graft copolymerization vinyl of ACKNOWLEDGMENTS monomers onto macromolecules having active an pendant groupcanbeachieved either by redox initiationThis reported research was supported by the National with a Ce(1V) ion or by photo-induced charge-transfer Natural Science Foundation of China. The authors are initiation with BP, dependingon the structure of the acgrateful to those coworkers whose research works have tive

Ceric Ion Redox Initiation Systems

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36 Modification of Properties of Nitrile Rubber Susmita Bhattacharjee Philipps Universitat, Marburg, Germany

Ani1 K. Bhowmick and Bhola Nath Avasthi Indian Instituteof Technology, Kharagpur, India

1.

INTRODUCTION

Nitrile rubber (NBR) was first commercialized by I.G. Farbindustry, Germany, in 1937, under the trade name of Buna N. Its excellent balance of properties confers it an important position in the elastomer series. Nitrile rubber, a copolymer of butadiene and acrylonitrile, is widely used as an oil-resistant rubber. The acrylonitrile content decides the ultimate properties of the elastomer. In spite of possessing a favorable combination of physical properties, there has been a continuous demand to improve the aging resistance of NBR due to thetougher requirements of industrial and automotive applications. Although, the heat resistance of NBR is directly related to the increase in acrylonitrilecontent (ACN)of the elastomer, the presence of double bond in the polymer backbone makes itsusceptible to heat, ozone, and light. Therefore, several strategies have been adopted to modify the nitrile rubber by physical and chemicalmethods in order to improve its properties and degradation behavior. The physical modificationinvolves the mechanical blending of NBR with other polymers or chemical ingredients to achieve the desired set of properties. The chemicalmodifications,on the other hand, include chemical reactions, which impart structural changes in the polymer chain. Bhattacharjee et al. [l] recently did an exhaustive review on properties and degradation of NBR. The improvement of physical properties of NBR by both chemicalandphysicalmethodshavebeenthoroughly described. Theyexamined the available literature on various modified NBRs and discussed in detail the refur-

bishment of properties in comparison with NBR. The present chapter focuses on one of the most important chemical modifications:the hydrogenation of NBR. Hydrogenated nitrile rubber (HNBR) is the latest high-performance elastomer that has emerged as an important product in the last decade. The principal driving force behind the developmentof HNBR isthe increasing performance demandedof elastomers by the automotive and oil drilling industries. Even the conventional oil-resistant rubbers, namely, nitrilerubber, chloroprene rubber, and chlorinatedpolyethylene are reaching their performance limits. Fluoroelastomers, which are the possible substituents in these applications, are very expensiveand possess processingdifficulties. Hence, HNBR has been developed to bridge the price performance gap between general purpose oil-resistant rubbers and fluoroelastomers. It has high heat and aging resistance, good resistance to swellingin technical fluids, including those containing aggressive additives, and outstanding wear resistance in extremely adverse conditions. HNBR has recently been marketed by Nippon Zeon Co., (Japan) under the tradename of ‘Zetpol’ [21; and Bayer A-G (Germany) andPolysar Rubber Co. (Canada) under the tradename of Therban [3]and Tornac [4], respectively. The double bond present in the diene part of the elastomer is generally more susceptible to thermal and oxidative degradation. The selective hydrogenation of olefinic unsaturation in NBR imparts significant improvements in resistance to degradation andother properties, such as permeability, resistance to ozone and chemicals, and property retention at high temperature. 555

Bhattacharjee et al.

556

This chapter mainly aims at describing the various methods and processes developed for hydrogenation of nitrile rubber. The characterization, physical properties, and application of hydrogenated nitrile rubber are also discussed. Another small section on hydroformylation of nitrile rubber has been included.

II. HYDROGENATION OF NITRILE RUBBER

Nitrile rubber hydrogenation is one of the most significant developments in the field of scientific and technological interest. Modification by hydrogenation of olefinic unsaturation in NBR offers an economic route for the synthesis of a new specialty elastomer, which is not accessible by standard polymerization techniques. The hydrogenation of various diene polymers has been reviewed previously [5-71. The recent advancement on the chemical modificationof polymers, with special reference to homogeneous catalytic hydrogenation, has been extensively discussed by McManus and Rempel [g]. McGrath et al. [9]have described the functionalization of polymers by metal-mediated processes. The survey of the available literature reveals that conversion of nitrile rubber to hydrogenated nitrile rubber has been achieved by both catalytic and noncatalytic methods [5-71. In catalytic hydrogenation, polymer is reacted with hydrogen in the presence of either a heterogeneous or a homogeneous catalyst. Noncatalytic hydrogenation is carried out in the presence of a suitable reducing agent. The important points to be considered in these reactions are the extent of conversion, side reactions, and selectivity. The recent citations suggest that the most popular method for noncatalytic hydrogenation is by diimide reduction using hydrazide reagents [5-71. The advantage of noncatalytic methods over catalytic

methods isthat a special hydrogenation apparatus is not required. Despite this advantage, there are a few side reactions, such as isomerization, attachment of hydrazide fragment to the polymer chain, depolymerization, and cyclization, during diimide reduction. However, under carefully controlled conditions, side reactions can be minimized [ 101. The studies on catalytic hydrogenation of nitrile rubber are much largerin number as compared with noncatalytic hydrogenation. The scope of the present chapter is to analyze all these techniques critically.The use of heterogeneousand homogeneous catalysts in selective reduction of olefinic bonds in the butadiene unitof nitrile rubber is summarizedseparately in the following sections. The initial reports of this vast body of literature on preparation of HNBR mainly covers the heterogeneous catalysis. But the more recent advancements are in the field of homogeneous catalysis, since it offers a better understandingof quantitative hydrogenation.

A. Thermodynamics of Nitrile Rubber Hydrogenation

The thermodynamic feasibility of a chemical reaction can be assessed before the experimental studies. Though the thermodynamic data alone are not sufficient to indicate the favorable conditions of the reactions, they definitely give some indicationsabout the probable range of temperature and pressure under which the reaction can be carried out. The estimation of heat of reaction from thermodynamical data gives the magnitude of the heat effects during the reaction. It also allows the calculation of the equilibrium constant from the standard free energies of the reacting materials. The expected maximum attainable yield of these products of the reaction can be estimated from the equilibrium constant.

557


Bhattacharjee et al. [1l] have calculated the thermodynamic parameters for hydrogenation of acrylonitrile-butadiene copolymer.

Table 2 Values of K; and A G" of NBR Hydrogenation

at Different Temperatures Temperature (K)

373 363 343 323 298

+

n'H2

11

*

k

m

InKE x AG" x (kJ/gmol)

.-

4.4 -11.1

4.0 -10.8

3.7

3.4

- 10.6 -10.4

Source: Ref. 11.

and AH" = AHo

+ AaT + (1/2) A b P + (1/3) AcT3

m

\

(4)

CN

where k is the number of cis-1,4-unit, 1 is the number of trans-lP-unit, m is the number of 1,2-unit, and n is the number of acrylonitrile unit in NBR and: (1)

n'=k+/+m

The nitrile rubber used for the experimental studies contained 40% acrylonitrile and the values of k, /, m, and n for the elastomer were calculated from spectroscopic studies as 359, 559, 215, and 770, respectively. The relationship used for the calculation of A G" at room temperature (298 K)is given by: AGO = AH" - TAS"

(2)

where A H" and A S" are the standard enthalpy change and standard entropy change of the reaction, respectively. The standard heats of formation and entropies for hydrogenation are available fromthe literature [12], A S" and A H" of the nitrile rubber and hydrogenated nitrile rubber were calculated from group contributions [121. AG" values at higher reaction temperatures were calculated according to the relationship: AGO = A H 0 - AaT In T - (112) A b P AcT3 - (1/6) IT

+

(3)

The coefficients a, b, and c for hydrogenation were obtained from the literature [l31 and those for nitrile and hydrogenated nitrilewere calculated from a group contribution method reported by Rihani and Doraiswami [14]. All the necessary data are listed in Table 1. The integration constant 1 and A H 0 have been calculated by incorporating the values of AG" and AH" at 298 K in Eqs. (3) and (4). The equilibrium constant at atmospheric pressure and various temperature has been calculated according to the relationship: (5)

AG" = - RTIn K;

The value of A G" at various reaction temperatures and corresponding K; values are reported in Table 2. It can be understood from Table 2 that the formation of hydrogenatednitrile rubber fromnitrile rubber is thermodynamically feasible. 111. HETEROGENEOUSCATALYTIC HYDROGENATION

This section discusses the processes for nitrile rubber hydrogenation developed duringpast years using a heterogeneous catalyst. The published reports suggest that mostly the palladium catalyst has been used for nitrile

Table 1 Thermodynamic Parameters at Room Temperature (298 K)

A H"

kJ/gmol Sample NBR 40 mol% ACN H2 HNBR 40 mol% ACN Source: Ref. 11.

3.3 -10.3

AS kJkmol)(K)

b x lo2

a

13593.02 578.90 - 36657.07 569.33 8.46 -96993.84

c x 104

- 818.03 0.10

2055.53 576.36

6.42

0.10 ,

14653.92

-7.81 X 10-4 - 858.16


558

Table 3 Heterogeneous Catalytic Hydrogenationof Nitrile Rubber

Degree of

rogenation SI. yst No.

Time

1. 2.

3. 4. PdC12lC 5. PdlSilica 6. PdlC 7.

Silica n

8.

PdC Pd-CdC Pd-AYC

Temperature Isobutyl-methyl ketone Acetone Acetone 50 Acetone

-

Acetone Isobutyl-methyl ketone

+ lithium

Pressure (MW

("C)

-

-

6.0

60

-

-

-

5.0

50

5

5.0 5.0

Pd/Ti02

-

47

5

83 95.3 92.2

4

50

4

-

-

-

-

5.0

50

4

50

25 4

98.5

70

27 4 28 2 29 6

96.8 93.0 87.8

Acetone

5.0

50 Acetone Tetrahydrofuran Acetone

5.0

-

( W

89.5 Iodine value = 31 84.0

acetate PdlCaCO39. PdlC 10. 11. PdlC 12. PdO 13.

16

(h)

-

5.0 5.0

rubberhydrogenation. However, heterogeneous catalysts based on different metals have been used for hydrogenation of many other diene polymers [15]. The selection of an optimum catalyst system is determined bythe requirement for the selectivity in hydrogenation. Heterogeneous catalytic hydrogenation of NBR is generally performed by stirring a solution of NBR with an insoluble catalyst under an atmosphere of hydrogen gas. The reduced product is usually isolated simply by filtration of the catalyst, followed by precipitation of the rubber and drying. The selection of the solvent and hydrogen pressure for a given hydrogenation isdependent on the catalyst chosen, its surface area, and onits inert support where it isdeposited. In many cases the catalyst can be recovered and reused. Nippon Zeon Co. has published many patents on heterogeneous catalytic hydrogenation ofNBR.Initially, they used 5% palladium on carbon as the catalyst for hydrogenation of NBR in iso-butyl-methyl ketone to obtain only47% conversion [16]. However, the product on compounding and vulcanizing showed better tensile strength, elongation, hardness, ozone cracking, and high-temperature oil resistance than the corresponding NBR. Later they changed the reaction condition and particle sizeof the carbon carrier to improve the degree of hydrogenation [17]. The reaction condition and catalysts used for heterogeneous catalytic hydrogenation of NBR [16-291 is arranged in Table 3. Supported palladium catalysts have been widely usedfor hydrogenation of conjugated diene polymerscontaining the nitrile group. Mostof the reactions were carried out at 5.0 MPa hydrogen pressure, 50°C temperature, and 4-6 h reac-

50

-

-

Ref.

17 18 19 20 21,22 23 24

26

tion time[ 16-29]. A high degree of hydrogenation could be obtained by changingthe active porous support and using other metals along with palladium. Researchers from Nippon Zeon Co. have used palladium and aluminum or palladium on carbon for preparing highlysaturated nitrile rubber [18,19]. For the similar reaction in the presence of only palladium on carbon, the extent of reaction was quite low [19].Palladium supported on silica showed highercatalytic activity during hydrogenation ofNBR and could easily be removed from hydrogenation mixtures by simple filtration [21,22]. The variation of average particle size of powdered silica support changed the degree of hydrogenation. Acetone isthe commonly favoredsolvent for these systems. Tetrahydrofuran and isobutyl-methyl ketone have been used in some reactions. It is important to note that allthe reactions mentioned in Table 3 proceed selectively withoutreduction of the nitrile group. Kubo et al. [24] hydrogenated carbon-carbon double bond in NBR in the presence of silica-supported palladium and lithium acetate promoter. Buding et al. [25] suggested that acrylonitrile-butadiene copolymers can be hydrogenated to high conversion by using palladium on calcium carbonate. The catalyst could be easilyrecovered from the reaction mixture and reused. Kubo et al. [27] developed another catalyst for hydrogenation of diene polymers that had high activity and could be easily recovered for reuse. It consisted of1% palladium supported on titanium dioxide and saturated 97% ofthe doublebonds present in nitrile rubber containing 35% acrylonitrile. Recently, Kubo [301 discussed the development of palladium catalyst on a special silica carrier

559


for selective hydrogenation ofNBR. He reviewed [31,32] the manufacture of hydrogenated NBRin the presence of supported palladium catalysts and its improved physical properties. Takahashi et al. [29] developed a polymer-supported catalyst for hydrogenation of conjugated diene polymer. They stirred a solution of NBR with the catalyst (PdO) at room temperature for 20 h, filtered the solution, and then conducted the hydrogenation. The pretreatment of the catalyst with the polymer enhanced its activity. It can be summarized from the available data in Table 3 that supported palladium catalysts selectively hydrogenated carbon-carbon double bondsin the presence of the nitrile group in NBR. However, there is no detailed fundamental study on heterogeneous catalytic hydrogenation of nitrile rubber in the literature that can provide an insight into the reaction. The available information is limitedsince most of the literature is patented.

depends on the cost of the metal and stability of the complex or ease of handling the complex. The hydrogenation procedure for most of the catalyst systems reported follows the same route. At first the polymer is dissolved in a suitable organic solvent. The solution is purged with hydrogen and the then requisite amount of catalyst is added. The mixture is subjected to the desired temperature and hydrogen pressure. After the completion of the reaction, the mixture is cooled andthe polymer isprecipitated out and finally dried. The reaction can be controlled up to the desired level of hydrogenation by proper combinations of the reaction pressure, temperature, time, and catalyst concentration. The homogeneous catalyst systems developed so far for NBR hydrogenation are discussed in the following sections. They have been divided into three broad classes based on the metals.

Rhodium Catalyst The major breakthrough in the history of selective hydrogenation is the discovery of the Wilkinson’s catalyst tris(tripheny1phosphine) chloro rhodium(1) [33],i.e. [RhC1(PPh3hl. This catalyst offers remarkable activity toward hydrogenationof carbon-carbon double bondin the presence of other reducible functional groups. Several other similar complexes have been introduced in the homogeneous catalyst series. Rhodium (Rh) complexes have beenextensivelyused for hydrogenation of carboncarbon double bond (=) in NBR. The major advantage offered by Rh complexes is their selectivity to hydrogenate C = C without causing any reduction of nitrile (CN) groups. The hydrogenation of CN groups, if any, can be considered negligible. the In presence of a nitrile group, which inhibits the catalytic activity during hydrogenation, Rh complexes are capable of retaining their high activity without any pronounced difference. The main criteria ofNBR hydrogenation is the selectivity toward reduction, in order to maintain the oil resistance and other physical properties of the hydrogenated product. It can be noted that possibly the most effectivecatalyst for selective hydrogenation of olefinic unsaturation in NBR is the Wilkinson’s catalyst [15]. There are quite a few reports of fundamentalstudies on this catalyst system. One of the reasons for the increasing interest in this complex is the simple method to synthesize this in high yield and the ease of handling due to its great air stability as a solid. McManus andRempel[8] have very systematically reviewed the commercial application of Wilkinson’s catalyst in preparation of HNBR. Generally, the desirable reaction condition for a commercial process to hydrogenate NBR using precious metal catalyst is the use of minimum catalyst concentration with respect to the polymer. Hence, an idealoperating condition in an industrial scale would involve temperatures greater than IOO’C, pressure greater than 2.7 MPa, and A.

IV.

HOMOGENEOUSCATALYTIC HYDROGENATION

In the last 15 years, there has beena tremendous explosion in the literature on the homogeneous catalytic hydrogenation of NBR. This section discusses the available reports on the development of various methods and processes for preparation of hydrogenated NBR in the presence of homogeneous catalysts. The principle of homogeneous catalytic hydrogenation is the activation of molecular hydrogen by a transition metal complex in solution andsubsequent hydrogen transfer to an unsaturated substrate. The most important advantage offered by homogeneous catalysts is their selectivity as compared with heterogeneous systems. The homogeneous catalysis also extends the opportunity to tailor ligands so as to enhance the reactivity and selectivity of the catalyst. Since the reaction occurs in the homogeneous phase, it also facilitates the study of the mechanistic steps involved in the catalytic cycle. However, it has the disadvantage that the catalyst remains in the product due to difficulty in removal.Even then, most of the previous literature suggests that the homogeneous catalytic methods for hydrogenation of unsaturated polymers are preferable to heterogeneous ones. The hydrogenation of functional unsaturated copolymers presents a special problem because many of the functional groups coordinate with the catalyst and poison its activity or are themselves reduced. These problems are particularly acute for highly coordinating functionalities such as nitrile, carboxyl, amino, hydroxyl, etc. Moreover, when the polymer contains two or three functional groups that are all prone to reduction, careful selectionof the catalyst system and reaction condition becomes important if regioselectivity during hydrogenation is desired. The choice of the catalyst also


560

Table 4 Rhodium Catalysts Used for NBR Hydrogenation

Degree of hydrogenation SI. (MP@ No. Solvent

Time Catalyst

Temperature

Pressure

1. 2.

fiCl(PPh3)~ RhCI(PPh3)3

Chlorobenzene Chlorobenzene

7.0 6.0

3.

RhCI(PPhd3

Chlorobenzene

4.

RhCI(PPh3)3 RhCI(PPh3)3 RhCl(PPh3)s RhCl(PPh3)3 RhCl(PPh3)s RhH(PPh3)d RhH(PPh3)d RhHz(dppb)2 (1,3-diphenyltriazenido) RhH(dibenzophosphole)3 Rh(C0D) (PPh3)2 OTs RhCl(sodium diphenylphosphino-benzenem-sulfonate)3

Chlorobenzene 2-Butanone Chlorobenzene Chlorobenzene Chlorobenzene Chlorobenzene Chlorobenzene

0.240-0.275 6.9 19.0 co.101 5.6 15.0 5.6 1.4-2.8 0.09

5. 6. 7. 8. 9. 10. 11. 12. 13. 14.

-

Water

polymer concentration around 2-10% by weight.On the basis of these criteria, Bhattacharjee et al. [l l] have optimized the reaction condition ofNBR hydrogenation, which mightbe considered usefulfor a commercial process. It is also necessary to add an excess of free triphenyl phosphine (PPh3)to the reaction mixture to achieve quantitative hydrogenation of C=C in the presence ofRhCl(PPh3)3. The choice of solvent is also important for efficient hydrogenation. The butadiene unit of NBR has cis-, trans-, and vinyl double bonds. The preferential reduction of these double bonds are different for various solvents used. Also the catalyst does not show its proper activity in the presence of many solvents. The continued interest in the use of the Wilkinson's catalyst incommercialapplicationand fundamental studies [11,34-471 has been chronologically organized in Table 4. The first patent on NBR hydrogenation, in the presence of RhCl(PPh& appeared in 1972 1341.Oppelt et al.[35] hydrogenated NBRin chlorobenzene solution using RhCI(PPh& to obtain a polymer with complete saturation of vinyl and trans double bonds. Weinstein [36] prepared elastomeric tetramethylene-ethyl ethylene acrylonitrile copolymer byselective hydrogenation of butadiene acrylonitrile rubber in the presence of RhCl(PPh3)3 andexcess of triphenyl phosphine. Almost complete hydrogenation wasobtained at both low- and high-hydrogenationpressures. However, at low-hydrogenation pressure appreciable excess of triphenyl phosphine wasrequired. Weinstein [36] observed that physi-

-

4.9 0.101

100 100 70-75 100 120 40 100 120 100 87- 130 55

-

100 75

Ref.

(h)

(%)

8 5.5

92 Iodine value = 14 99.7

34

99.7 100 100 99.2 93(L-XNBR) 51-99.95 91

37 38 11 39

-

-

10

99.5 60

44 45 46

12 5 11

-

11 2-15 19

-

-

-

35 36

40 41

42 43

cal and chemical properties of modified polymer gums and their sulfur and peroxide cured vulcanizates were resistant to attack by oxygen and showed unexceptionably high tensile strength, which was attributed to the crystalizability of the gumson stretching. Peroxidecured polymer vulcanizates prepared from green stock containing magnesium oxide had a combination of resistance to heat aging inair, hydrogen sulfide, and oil, suggesting their applicability to oil drilling components. The industrial applicationof RhCl(PPh3)3 has stimulated a few academic studies to obtain a better understanding of hydrogenation of NBRsubstrate. Thermodynamics and kinetics of this reaction have been thoroughly elucidated. This has given valuable insight into the possible mechanism for the hydrogenation process. Mohammadi and Rempel [38], for the first time, conducted a detailed study involvingthe hydrogenation of acrylonitrile-butadiene copolymer inthe presence of RhCl(PPh&. Under mild reaction conditions, this catalyst provided quantitative hydrogenation of carbon-carbon unsaturation without any reduction of the nitrile functionality. The reactions were carried out in a ketone solvent (2-butanone)at 40°C under a hydrogen pressure < 101.32 KPa. The selectivity of the catalyst for terminal versus internal double bondspresent in the polymer was markedly influenced by the nature of the solvent. In 2butanone, the catalyst did not exhibit selectivity for the hydrogenation of 1,2- over 1,Cunits of the copolymer, whereas a distinct preference toward saturation of the 1,Zunit was observed when the reaction medium was


chlorobenzene. They also observed that, under this mild condition, there is no need to add an excess of PPh3 to maintain the catalytic activity. During the hydrogenation, an excess of PPh3prevents the formation ofa binuclear Rh complex that is inactive as a hydrogenation catalyst. Under relatively mild conditions, the nitrile group of the polymer might inhibit the formation of inactive catalytic species, whereas the contrast is observed at higher reactiontemperature and pressure [ 111. At elevated reaction conditions, it is necessary to add an excess of PPh3 to stabilize the rhodium complex and prevent the reduction of the complex to Rh metal. The reaction kinetics showed significanteffects on the concentration of olefinic substrate and hydrogengas as well as solvent. A first order dependence of the hydrogenation rate on the catalyst concentration wasobtained. The reaction was first order at low-hydrogen pressure and zero order at higher hydrogen pressure. The proposed mechanism based on the kinetic studies and spectral observation is illustrated in Scheme 1. Mohammadi and Rempel[38] suggested that the active catalyst species, RhClH2(PPh3)2 interacted with C=C in the rate-determining step. The nitrile functionality present in the copolymer also reacted with the active catalyst and inhibited the rate of hydrogenation to some extent. The effectof temperature on the reaction rate provided the activation parameters associated with the rate constants. The activation energy of the reaction had a value of 87.28 kJ/mol and at 40°C the activation enthalpy and activation entropy of the reaction were estimated to be 84.67 kJ/mol and -8.3 J/(mol K), respectively. Bhattacharjee et al. [l11 optimized the reaction conditions for hydrogenation ofNBRin the presence of RhCl(PPh3h. In order to identify the important factors

561

that influenced the nature and extent of the reaction, parameters such as temperature, pressure, time, catalyst concentration, and solvent were vaned. When the reaction wascarried out in chlorobenzene, there was no conversion at room temperature. The extent of the reaction increasedwith an increase in temperature, hydrogen pressure, and catalyst concentration. The authors found that the reaction of NBR at a catalyst concentration of 0.02 mmol, under 5.6 MPa hydrogen pressure, at 100°C in chlorobenzene for 11 h was optimum for complete hydrogenation. The amount of catalyst required under this condition is much lower as compared to that reported by Mohammadi and Rempel[38]. This supports the practical importance of the reaction optimized at the elevated condition. Bhattacharjee et al. [l11 also attempted to study the kinetics of the NBR hydrogenation catalyzed by RhCl(PPh3)3 at higher reaction pressure and temperature. The experiments were conducted at various intervals of time and corresponding hydrogenated polymer samples werecharacterized to know the extent of reaction. The experimental data were examined to understand the dependence of reaction variables on the rate of hydrogenation. The observations made were the same as that reported earlier by Mohammadi and Rempel[38]. The activationenergycalculated for the reaction at 100°C was 22 kJ/mol. The value of activation energy is significantly lower than that reported by Mohammadi and Rempel [381. The formation of the active catalyst precursor R ~ C I H Z ( P Pmight ~ ~ ) ~be more favorable at elevated temperatures. Bhattacharjee et al. [ 1l] characterized the hydrogenated product and studied its physical properties. They found that no degradation or side reaction occurred during the hydrogenation process and

- H2 RhCl(PPh3)2

/

(fast)

Scheme 1 Mechanism of NBR hydrogenation catalyzed by RhCl(PPh3)3. Source: Ref. 8.

562

the molecular weights of the polymer did not change significantly. RhH(PPh3)4has been also used for NBR hydrogenation [41,42]. It is claimed that the analogous complexes ofRhCl(PPh3)3 containing5-phenyl-5Hphenylphosphole ligand [44], RhHz(dppb)z( 1,3-diphenyltriazenido) [43] and Rh(COD)(PPh3)2OTs [45] offer excellent activity for NBR hydrogenation. Recent patents have described the hydrogenation of low-molecular weight NBR [47] and butadiene-isoprene-(meth)acrylonitrile copolymers [39] using RhCl(PPh3)s. The rhodium complexes are excellent catalysts for hydrogenation of NBR. At low temperature and pressure, high catalyst concentrations are used to obtain a better rate of reactions. Due to higher selectivity of the reaction, pressure and temperature can be increased to very high values. Consequently the rhodium concentration can be greatlyreduced, which leads to high turnover rates. The only practical drawback of Rh complex isits high cost, This has initiated the development of techniques for catalyst removal and recovery (see Section VII), as well as alternate catalyst systems basedon cheaper noble metals, such as ruthenium or palladium (see Sections 1V.A and B).


attributed to the fact that the formation of the active catalytic species: RuHCI(PPh3)3 is promoted by these additives. The nonketone solvents can be used for certain Ru complexes containing organic phosphine ligands when RC02H and PPh3 are the additives [57,74]. Rempel et al. [67,68] have reported that the complexes of the form RuXY(CO)L2Z (X and Y may be a variety of anionic ligands, L is a bulky phosphine such as tricyclohexyl or triisopropyl phosphine, and Z may be anyneutral coordinating ligand) offer quantitative hydrogenation ofC=C inNBRin nonketone solvents without anygel formation. Besides the cost factor, these catalyst systems are also more practical in commercial fields due to higheroperating temperature and pressure. Quantitative reduction of olefinic unsaturation in NBR can be achieved at elevated temperature and pressure without addition of any excess ligands (such as PPh,) for maintaining the catalytic activity. However, these complexes cause crosslinking of the elastomer during hydrogenation. This was observed from the unusually high viscosities of the resultant hydrogenated polymer [61]. Furthermore, it was also found that the Ru complexes catalyzed little reduction of nitrile groups in the elastomer. Thismight be a reason for the high viscosities of the polymer. It was later found that the nitrile group was converted to secondary amines. This could be preB. RutheniumComplexes vented by the addition of small amounts of primary amines to the reaction system [61]. They might react with the imines, whichare the primary product of nitrile Hydrogenation of olefinic unsaturation using ruthenium hydrogenation, thus preventing the crosslinking. The (Ru) catalyst is well known.It has been widely usedfor aqueous solutions of carboxylicor mineral acidor transiNBR hydrogenation. Variouscomplexes of Ru has been tion metal salts can inhibit crosslinking caused by the developed as a practical alternative ofRh complexes ruthenium catalyst [73,74]. They probably avert crosssince the cost ofRu is one-thirtieth of Rh. However, linking by inducing hydrolysis of imine groups produced they are slightly inferior inactivity and selectivity when by hydrogenation of nitriles. The crosslinking and gel compared with Rh catalyst. formation are the major drawbacks of the Ru catalyzed Reports on NBR hydrogenation catalyzed byRu hydrogenation process. Consequently, many answers complexes [48-741 are summarized in Table 5. A variety have arisen to surmount this problem. ofRu complexes, such as RuCl2(PPh3)3[48-521, RuRecently, Guo et al. [67,73] havestudied the hydroH(OZCR)(PP~~)~ [54], R u H C ~ ( C O ) ( P P ~[57], ~ )and ~ genation ofNBR in the presence of RuCI(PhC02)RUC~(O~CR)(CO)(PP~~)~ [58] havebeenused as the (CO)(PPh3h.They conducted the reaction at relatively catalysts for hydrogenation of NBR. The easy accessilow temperatures (65-120°C) and pressures (approxibility of these complexes isa great advantage in commermately 0.1 MPa) andobtained only 40% hydrogenation. cialization of the process. Almostall the ruthenThe activity of this catalyst was poorer than ium-phosphine complexes offer adequate activity for RhCI(PPh3)3,under similar conditions. This complex reduction of carbon-carbondoublebonds inNBR. also catalyzes the isomerization of C S in the polymer However, the main drawback of these systems is the gel backbone. Dueto positional isomerization,C S shifted formation duringthe hydrogenation process, unless the conjugated to nitrile groupsof the elastomer. Such C = C reaction is carried out in ketone solvents. are difficult to hydrogenate as they are stabilized byelecTable 5 indicates that R u c l ~ ( P P h ~ has ) ~been fretron delocalization due to conjugation. Later, they could quently used for selective hydrogenation ofC=Cin achieve more than 97% hydrogenation, at low-polymer NBR [48-521. This is commercially available andalso is concentration, using higher pressure and temperature. easy to synthesize. In most of the patented processes, The gel formation in this case was also less. low-molecular weightketone solvents are used to avoid Martin et al. [69] undertook a study of the kinetics the gel formation. The activity of the catalyst can be and mechanism of NBR hydrogenation using various Ru enhanced by the use of certain additives, such as trieththe activity ofRuXCIylamine [59],isopropanol [52], and ammonium hexaflu- complexes.Theyexamined (CO)L2(X = H, Ph, or CH=CHPh; L = PCy3, PiPr3, orophosphate [50] in the reaction system. This might be

on


563

Table 5 Hydrogenation of NBR using Ruthenium Complex Time No. Solvent

Temperature

SI.

1. 2. 3. 4. 5. 6. 7. 8.

9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21.

22. 23. 24.

Catalyst

Pressure (h) ( M W

Butanone RuC12(PPh3)3 RuC12(PPh3)3 RuCIz(PPh3)3 R~Clz(PPh3)3 Acetone RuC12(PPh3)3 2-Butanone RUCI~(PP~~)~/NH~PF~/E~~N RuH(MeCOz)(PPh3) Butanone RuX[(LI)(L2)n],X = halo, SnCI2, Low molecular H L, = halo, hydrido, indenyl, weight fluoroenyl, L2 = phosphine, ketone bisphosphine or arsene, (n = 1 or 2) 2-Butanone RuH4(PPh3)3 RuHCI(CO)(PPh& + CHaCOOH Chlorobenzene RuH(OAc)(PPh& Acetone RuCIz(PPh3)3 + Et3N Acetone RuCI((C0) (tricyclohexylChlorobenzene phosphine)2] RuC13.3HtO + PPh3 + EbNCI Acetone RuClp.3HzO + PPh3 Acetone RuClp.3HzO + PPh3 + Acetone CaCI2.2H20 RuCI3.3H20 + PPh3 Acetone RuCl3.3H20 + PPh3 + AcOH Acetone RuCl [(CO)(COPh) (PPhdzI Ru(x)CI(Cb)L2, X = H or styryl L = tricyclohexyl or triisopfopyl phosphine RuXY(CO)ZL2or RuX(NO)(CO)L2 Chlorobenzene X = ,halogen, carboxylate or phenyl vinyl group, Y = halogen, H, Ph, phenyl vinyl, carboxylate, Z = CO, pyridine, PhCN, (MeO)3P, L = Hindered alkyl phosphine RuCIz(styry1) (tricyclohexyl 2-Butanone phosphine)z RuC12(PPh3)3/FeS04/(NH,)S04 2-Butanone Ru(I1) complex + HOAc Organic solvent + water

or PPh3) toward selective hydrogenation of NBR. They observed that the Ru complexes withPPh3 ligandcause gel formation, whereas the complexes with PCy3 and P'Pr3 ligands steer efficient hydrogenation without gel formation. They also concluded that the gel formation in the Ru catalyzed hydrogenation of NBR innonketone solvents is onlydue to reduction of nitrile groups during the reaction process. The kinetics of hydrogenation using RuCl(CH=CHPh)(CO)(PCy3)was also examined byMartin et al. [69]. The kinetics of hydrogenation showed afirst order dependence on C = C concentration,

14.0 4.9 -

4.9 14.0 4.9 14.0

-

14.0 1.S

-

5.6 9.8 9.8

Degree of ("C)

(%l

150 145 145 145 135

99. l 93 100 95.9 99.8

145

99

-

-

Ref. 48 49 50 51 52 53 54 55

130 140 145 145 140

4 1.8 6 4 2

90.0

145 145 145

5 4 3

94.5 92.2 85

64

145 155

-

94 97.2

66

6

95 99.5 92.0 99

56 57 58 59 60, 61

62 63

65 67.68 69

5.0

140

0.5

99

-

-

71 -

99

-

-

72,7374

99 99.8

-

-

70

hydrogen pressure, and catalyst concentration and an inverse dependence on nitrile concentration. They postulated a mechanism (Scheme2) for hydrogenation using this catalyst. The coordination of the nitrile group to the metal center is the reason for hydrogenation of these groups. The mechanism is slightlydifferent from the Rh catalyst where no ligand is displaced from the metal center. Table 5 lists other ruthenium complexes that could catalyze selective hydrogenationinNBR. However, their activity could not be properly compared as they

564

Bhattacharjee et ai.

NCR

I

RCHZCHZR (hydrogenated Polymer)

RuHCI(CO)(PC~&

Scheme 2 Mechanism of NBR hydrogenation catalyzed by RuCl(CH=CHPh)(CO)(PCy,),.

occur under different reaction conditions. In spite of possessing certain disadvantages, the ruthenium catalyst is anattractive alternative for selective and quantitative hydrogenation of nitrile rubber. Guo et al. [70,71,73] recently attempted to hydrogenateNBRinemulsionformusingRu-PCy3complexes. However, successful hydrogenationcan only be obtained whenthe emulsion is dissolved in a ketone solvent(2-butanone).AvarietyofRu-phosphinecomplexes have been studied. Crosslinking of the polymer could not be avoided during the reaction. The use of carboxylic acids or first row transition metal salts as additivesminimized the gel formation. The reactions under these conditions require a veryhigh catalyst concentration for a desirable rate of hydrogenation.

'

Source: Ref. 8.

C. PalladiumCatalysts In the preceding section, it has been shownthat considerable attention has been devoted to palladium as a heterogeneous catalyst. The present section describes the homogeneous palladium catalysts developed for hydrogenation of NBR. The main drive behind the development of various catalyst systems isto find suitable substituents of the Rh catalyst. Palladium complexes are much cheaper as compared with Rhand exhibit comparable activity and selectivity to Rh and Ru complexes. The most widely used homogeneous palladium catalyst is palladium carboxylate. Table 6 lists the palladiumbased catalyst systems developed for NBR hydrogenation [75-871. They are regarded as homogeneous cata-

Table 6 Hydrogenation of Nitrile Rubber in Presence of Palladium Catalyst

Degree of SI. No.

1. 2. 3. 4. 5.

6. 7. 8. 9. 10. 11. 12. 13.

Solvent Pd[(OAc).rls Pd[(OAc)2]3 reduced with hydrazine Pd[(RC02)2]3 R = Me or Pr Pd[(OAc)zl3 [Pd(OAc)Q-benzoyl pyridine)]^ [Pd(OAc)(2-benzoylpyridine)]^ [PdCh(MeCN)21 PdO(H20),/MeCN/PhP04 PdO complexed with NBR Pd(2-ethylhexanoate)/methyl aluminoxane Pd(2-ethylhexanoate)IEt3AI Pd(ben~oate)~ Pdt(OAc)tl3~NO3

Pressure Temperature Time hydrogenation (MW ("C) (h)

( W

97.2 98.5

4.9 4.9

50

3 5

-

-

-

-

-

Acetone Acetone Acetone Acetone Tetrahydrofuran

2.7 5.6 2.7 4.9 2.41

60 75 60

1 6 1 6

96 80 68 84.9 70.0

Acetone Acetone

-

-

50

75-85

-

2-Butanone Benzene Water

6.2 3.0 3.0

60 50 50

2 6 6

75 76

84

77 78 79 40,80 81 82 83 84

90 90.2 78.2

85 86 87

-

2-Butanone

Ref.


lysts because they are dissolved inthe polymer solution during hydrogenation. It is presumed that they are reduced to colloidal Pd metal under the hydrogenation condition. In a patent, Kubo et al. [75] mentioned that the catalyst can be removed fromthe polymer solution by filtration or centrifugation. This might be attributed to the formation of a uniform dispersion of very finely divided metal throughout the substrate solution when the catalyst is subjected to the reaction. The palladium complex should be properly soluble in the reaction mixture in order to incur a high degree of hydrogenation. This is supportedby the studies on PdCI2catalyzed hydrogenation of NBR [75]. PdClz is not a readily soluble complex inthe polymer solution and, hence, offers inferior levels of conversion as compared to palladium carboxylates. Acetone isthe best solvent for NBR hydrogenation in the presence of palladium carboxylates. No hydrogenation is achieved when chloroform or chlorobenzene are the solvents. Since it is understood that palladium is reducedto colloidal metalin the presence of hydrogen, attempts have also been made to reduce the palladium by hydrazine [76], methylaluminoxane[841, and trialkyl aluminum [85] to improve the catalytic activity. Bhattacharjee et al. L781 have conducted a fundamental study on the catalytic activity offered by palladium acetate toward hydrogenationof C=C in nitrile rubber. Palladium acetate is a simple, commercially available, air stable, and relatively less expensive complex of palladium. It exhibited excellent selectivity and hydrogenationefficiency.With increases in reaction time, pressure, temperature, and catalyst concentration, the extent of hydrogenation increased. A maximum conversion of 96%could be achieved at 60°C under 2.7 MPa pressure and 0.54 mmol catalyst for .l h in acetone. The reaction is very rapid even at room temperature. From the spectroscopic analysis, they confirmed that the CN group was not hydrogenated during the reaction. A kinetic study of NBR hydrogenation suggestsa first order dependence on the olefinic substrate. The rate of hydrogenation increases with increases in catalyst concentration and hydrogen pressure. The activation energy of the reaction calculated from the Arrhenius plot is 29.9 kJ/mol. This is comparable to that obtained for the Rh catalyst [ 1l]. The apparent activation enthalpy and activation entropy are 27.42 kJ/mol and -0.20 kJ/(mol-K). Gel permeation chromatographic measurements show that there is no change in the molecular weight of the polymer after hydrogenation. The glass-transition temperature reduces gradually withthe increase in the level of hydrogenation. They also studied some of the physical properties of the hydrogenated product. Bhattacharjee et al.[79] introduced another new catalyst based on a Pd complex containing bothacetate and benzoyl pyridine ligands (Table 6). This was developed to hydrogenate liquid carboxylated nitrile rubber (L-XNBR) [go].Selective hydrogenation of C=C in L-

565

XNBR is moreintricate due to the presence of two additional functional groups: carboxyl andnitrile. Palladium acetate does not hydrogenate L-XNBR, whereas cyclopalladate complexof 2-benzoyl pyridineeffects measurable amounts of hydrogenation [40]. The reduction was strictly restricted to C S in the presence of carboxyl andnitrile groups. RhCI(PPh&, on the other hand, cause significantamounts of decarboxylation along with hydrogenation.Hydrogenation ofsolidnitrile rubber with different acrylonitrile content and solid carboxylated nitrile rubber (XNBR) hasalso been accomplished by using this catalyst [791. The physical properties of

II

Scheme 3 Mechanismofhydrogenation

of NBR catalyzed by Pd(OAc)(Zbenzoyl pyridine)]*.Source: Ref. 79.

566

the modified rubbers were also examined. In contrast to palladium acetate, chloroform andchlorobenzene could also be used as solvents. The authors have proposed a mechanism for hydrogenation in the presence of cyclopalladate complex (Scheme 3), which is similar to that reported for Rh and Ru catalyst (Schemes 1 and 2). The first attempt to hydrogenate NBR in emulsion form wascarried out in the presence of palladium benzoate [86]. However, an organic solvent and high concentration of catalyst were required to attain a reasonable level of hydrogenation. A similar observation was experienced later for Ru catalyzed hydrogenation in emulsion form. A very recent patent from Nippon Zeon Co. reveals that NBR latex can be hydrogenatedby using palladium acetate in the presence of nitric acid [871. This is a very significant development as it does not require any organic solvent. The available data in Table 6 reveal that palladium complexes are excellent catalysts for selective hydrogenation of C=C in NBR. Recent attempts to recover the catalyst (see Section VII) after hydrogenation and lower the cost of the metal make it an attractive supplement in the industrial production of HNBR. V. HYDROGENATION OF LIQUID RUBBER The liquid nitrilerubbers are generally usedas nonvolatile and nonextractable plasticizers. They also function as binders and modifiersfor epoxy resins. Their moderate heat resistance limits their ability to meet industrial requirements. Hence, attempts have been made to improve their thermal andoxidative resistance by saturating the polymer backbone. Hashimoto et al. [88] have hydrogenated liquid nitrile rubbers for the betterment of the thermo-oxidative resistance and retention of properties at higher temperature. Studies onhydrogenation of liquid rubbers are quite challenging, as most of them contain more than one reducible functionality. Hydrogenation of liquidcarboxylatednitrile rubber (L-XNBR)complicates the method, because boththe carboxyl and nitrilefunctionality interfere in the double bond reduction. Palladium acetate, which evolved as an efficient catalyst system for selective hydrogenation of solid NBR, could not hydrogenate L-XNBR. Wilkinson’s catalyst offered reduction of C 4 along with a significant amount of decarboxylation and, under certain conditions, hydrogenation of nitrile groups [40]. Bhattacharjee et al. [40,80] have developed a new catalyst based on the palladium complex for quantitative hydrogenation of C=C in LXNBR. Successfulhydrogenation of C=C was obtained without disturbingthe carboxyl and nitrilecontent of the polymer. These groups also did not restrict the catalytic activity of the complex. The degree of hydrogenation at a particular reaction condition depended mainly on the


ratio of catalyst to that of the diene unit of the liquid rubber. The reaction kinetics exhibita first order dependence with respect to the olefinic substrate [go]. This is in line with the earlier observation on Rh [l 1,381 and palladium acetate [78] catalysts. An increase in the amount of nitrile andcarboxyl groups decreased the rate constants of the reaction. The activation energy of the L-XNBRhydrogenationwas 20.2 kJ/mol. The glasstransition temperature of the saturated liquid polymer decreased significantly. VI. HYDROGENATION OF NITRILE RUBBERLATEX The idea of hydrogenating nitrilerubber in latex form is of high practical importance in the commercial production of HNBR. This could lead to a direct preparation of HNBR fromthe monomers without isolationof NBR if hydrogenation could be introduced at the emulsion stage. The expense of using an organic solvent and subsequent removal couldalso be curtailed. Attempts have beenmade to hydrogenate NBR emulsionusing Ru [71-731 and Pd [S61 catalysts. However, they all require an organic solvent to swell the emulsion anda high concentration of catalyst for a sufficient degree of hydrogenation. Recently there was a report of NBR latex hydrogenation in the presence of palladium acetate which does not require any additional use of organic solvent [87]. Researchers from Nippon Zeon Co., have developed a novel method for selective hydrogenation of C = C in ‘NBR latex by using palladium acetate and nitric acid. The introduction of an equimolar amount of nitric acid along with palladiumacetate helps to avoid the organic solvent, which is normally usedto dissolve the catalyst and swell the latex particles. However, the investigations in this area are not sufficient to generalize the method. Singhaet al. [46] developed a water-soluble analog of Wilkinson’s catalyst for hydrogenation of NBR in the latex stage.RhCl(sodiumdiphenyl phosphino benzene-m-sulfonate)3 effecteda moderate degree of hydrogenation of C = C in NBR. However, hydrogenation was accompanied by an increased gel content of the latex. The noncatalytic methods have also been usedfor latex stage hydrogenation of NBR (see Section VIII). The research in this area has a great potential from scientific and technologicalaspects and requires further exploration. However, the reported attempts are a welcoming endeavor to hydrogenate nitrile rubber in latex form. VII. REMOVAL OF THE CATALYST The only weakness of the homogeneous catalytic hydrogenation of elastomer is the removal of the catalyst from

of

567

Modification

the polymer mixture. This is the main reason for the high cost of hydrogenated nitrilerubber. If the expensive catalyst can be recovered and reused, the difference in the cost of NBR and HNBR can be brought down to a minimum level. The service life and resistance to hightemperature degradationmightbe enhanced further when these active metals are removed fromthe polymer substrate. Several researchers have come up with new ideas to remove the catalyst from the hydrogenated polymer. Ahlberg et al. [89] reported the removal of RhCI(PPh3)3 catalyst residues from HNBR. They kneaded the rubber with 20-500 phr methanol in a mixing apparatus at 60°C. The rubber was later separated from methanol.The polymer sample, which initially contained 116 ppm Rh and 1.46 wt% PPh3, after extraction showed 86 ppm Rh and 1.18 wt% PPh3. Osmanet al. [90] developed another procedure for removal of Rhresidues from HNBR solution.The reaction mixture (200 mL solution of chlorobenzene containing 6 g of HNBR), after catalytic hydrogenation, was treated with 10 g of starch solution in water, refluxed at 110-120°C for 6-7 h, and cooled to ambient temperature. A precipitate, formed upon cooling, was separated by filtration. The rubber was coagulated with methanol and dried. The residual Rh content of the elastomer decreased to 5 ppm from 85 ppm. Madgavkar et al. [91] loweredthe level of catalyst residues below 5 ppm in polymer solutions. They oxidized the polymer solution in the presence of oxygen and then treated with carbon black to remove the catalyst. Panster et al. [92] reported the recovery of Rh and Ru catalysts from solutions of HNBR. Morethan 80% of the catalyst could be recovered by absorption on siloxanes containing tertiary amino groups, sulphide, phosphine, thiourea, or urea. However, the research is still in progress and probably a more convenient method will soon be developed to completely recover and reuse the expensive catalysts.

VIII. NONCATALYTIC HYDROGENATION The available studies indicate that diimide has been used as a reducing agentfor the preparation of HNBR. It has been used mainlyas an alternative for hydrogenation of nitrile rubber latex. The use of diimide to hydrogenate low-molecular weight olefines is well known in the organic literature [93]. Diimide can be conveniently generated in situ by thermal treatment of solutions of p-toluenesulfonyl hydrazide or oxidation of hydrazine. Wideman [94] developed a method for reducing the C=C of NBR in the latex stage. The process involved refluxing NBRlatex with hydrazine, an oxidizing agent, and a metal ion activator. The latex was later filtered and coagulated. During the in situ generationof diimide,

there is formationof some by-products. The reaction of these by-products with unsaturated sites of the polymer leads to molecular weightdegradation. Apart from this, hydrazine can remain attached to the polymer chain. A novel process for preparation of highly saturated nitrile rubber latex has been introduced by Parker et al. [IO]. It involves the diimide reductionthat is generated in situ within the latex by the oxidation of hydrazine hydrate in the presence of a catalyst. The reaction operates at ambient pressure and at 40-70°C temperature. Nitrogen and water are the by-products of the reaction. The process is quite efficient for selective hydrogenation of C S in NBR and has a high commercial interest because it converts NBR latex into HNBR latex in a onestep reaction. The HNBR latex produced by this method has been usedin many latex-related applications. Scientific studies ondiimide reduction of low-molecular weight polymers by p-toluene sulfonyl hydrazide are also reported [951. Schiessl [961 attempted to scavenge hydrazine attached to the latex following the reduction procedure. NBR latex was reduced with hydrazine peroxide and at a 60-70°C temperature for 1-5 h in the presence of copper sulfate or ferrous sulfate initiators. The hydrogenated polymer latex was then treated with compounds consisting of various functional groups, such as isocyanates, acrylates, acids, ketones, diketones, or aldehydes in large excess to eliminate the free hydrazine in the system. In another process to minimize residual hydrazine in HNBR latex [97], they treated HNBR latex containing 3.76% hydrazine with oxygen at 0.28 MPa in the presence of 0.2% hydroquinone at 50°C for 20 h and lowered the hydrazine level significantlyto 100 ppm. Noncatalytic hydrogenationispromisingfrom a commercial point of view if the side reactions during the process can be carefully excluded. Moreover, this is of particular importance for online reduction ofNBR emulsion.

IX. HYDROFORMYLATION OF NITRILE RUBBER Hydroformylation of nitrile rubber is another chemical modification that can incorporate a reactive aldehyde group into the diene part and further open up new synthetic routes to the formation of novel nitrileelastomers with a saturated backbone containingcarboxyl or hydroxyl functionalities. Bhattacharjee et al. 1981 reported that nitrile rubber can be selectively hydroformylated in the presence of RhH(CO)(PPh3)3 and RhCl(CO)(PPh& under high pressure of carbon monoxideandhydrogen. They found that RhH(CO)(PPh&wasmoreefficient than RhCI(CO)(PPh3)2.However, both catalysts offered se-

568

Bhattacharjee et ai.

1

CO / H,

CATALYST

r

1

r

:H-CH2

l

t

,H “ C H 0

lectivity towardthe addition of aldehyde group to C=C in the presence of nitrile functionality; 30% conversion couldbeachieved at 90°C under 5.6 MPa pressure (CO:H2 = 1:1) in presence of a 0.43 mmol/L catalyst. When the hydroformylation was more than 30%, gel formation occurred. The rate constant was reduced by the higher partial pressure of carbon monoxide. With the increase in the acrylonitrileconcentration of the copolymer, there was a significant decrease in the extent of hydroformylation as well as the rate constants of the reaction under certain reaction conditions. Characterization of the products by spectroscopic techniques indicated the formation of internal aldehyde from a 1,4-unit and terminalaldehyde from a 1,Zunit of the diene fraction of the copolymer. The intrinsic viscosity of the modified polymer decreased with an increase in aldehyde incorporation. The glass-transition temperature increased with the extent of hydroformylation. The thermal stability of the hydroformylated product was, however, lower than that of nitrile rubber. Hydroformylated nitrile rubbers might have potential as binders for coupling withother rubbers due to the presence of polar aldehyde group. This might also open up the still greater potential of NBRcontaining saturated backbone in composite materials [l].

X. CHARACTERIZATION OF MODIFIED NITRILE RUBBER Characterization and understanding of the microstructure become important after hydrogenation and hydroformylation of the nitrile rubber since the amount and distribution of the residual double bonds influence the properties of modifiedrubber. The conventional analytical tools have been usedto characterize the elastomers. Spectroscopy is the most useful technique for determination of the degree of hydrogenation in nitrile rubber. Back [99,1001 described the infrared (IR) spectrophotometric methods for determination of the proportion of acrylonitrile, butadiene, and hydrogenated butadiene in hydrogenated acrylonitrile-butadiene rubbers of unknown composition. In addition, it wasalso demonstrated that the cisltruns ratio of the 1,Cbutadiene structure in NBR and the residual double bonds in HNBR could be calculated from the IR measurements without a knowledge ofthe correspondingabsorption coefficients. These calculations are based on the fact that the nitrile group isnot reduced during the hydrogenation process. This can be supported by the IR spectrum of a HNBR sample, which does not show any additional peaks due to hydrogenation of nitrile functionality. The selectivity


569

Wavenumber. cm-'

Figure 1 IR spectra of:1-NBRand2-5-HNBRwith58, 71,and99.9mol% hydrogenation, respectively. Source: Ref.11.

of the catalyst can also be established from the IR studies. Figure 1 illustrates the IR spectra ofNBR and HNBR samples with different concentrations of residual double bonds [l l]. The CN stretching vibration is observed at 2222 cm" in NBR and HNBR. The peak at 1440 cm" is for C-H deformation of "CHI groups. The =C-H out of plane deformations of trans, vinyl, and cis double bonds are observed at 970cm",920 cm", and 730 cm", respectively. These peak absorbances decrease gradually, and a new peakat 723 cm" appears on the spectrum of HNBR for " C H r rocking vibration [when(CH&, n > 41 [l 1,781. The CN stretching vibration is taken as an internal standard and the

h

Figure 2 'HNMR with 71 and 99.9 mol%

Source: Ref.11.

ratio of absorbances at 970 cm", 920 cm", and 730 cm" with respect to 2222 cm", is compared with the peak absorbance ratio at 723 cm" for an estimation of the level of hydrogenation. The degree of hydroformylation [98]was calculated from the relative decrease in the absorbance ratio due to =C-H and increase in the absorbance ratio at 723 cm". Marshall et al. [l011 developeda method for quantification of the level of hydrogenation in HNBR by nuclear magnetic resonance spectroscopy (NMR). The measurements are based onrelative amounts of olefinic and aliphatic signalsin proton NMR ('H NMR spectra). Figure 2 shows the 'H NMR spectra of NBR and two HNBR samples. The signals due to olefinic protons appear as a broad peak between 5.0 to 5.8 ppm. All the protons in the "CH3, l00 3.1 1.1 matrix. Nevertheless, as in our earlier studies W], the Twin-screw in situ fibrillation was more pronounced for the more blending (Fig. 3) viscous PP (VB1950K) and this PP grade was selected 34 4.1 3.1 1.1 Melt blend (L) 1550 for the additional experiments. 54 2.0 1.3 6.1 Melt blend (H) 3630 The morphologies of the PP-LCP melt blends preSubsequent pared by a twin-screw extruder at 290°C with low (L) single-screw and high (H) drawing speed are presented in Fig. 3 and extrusion (Fig. 4) the corresponding mechanicalproperties in Table 2. At He 1 25 1310 6.0 4.2 0.6 the lower draw ratio a clear skidcore morphology was 3.6 4.4 26 1570 He2 0.8 generated, whereas the higher take-up speed resulted 20there 1200 He3 0.78.0 3.8 in an extremely highly fibrillated morphology and 19 Es- 1160 He4 >l0 0.7 3.7 were virtually no differences between skin core. and He5 (less 1810 34 6.4 3.3 0.9 pecially withthe higher drawratio, it was expected that drawn) the fibers in the strand might beseveral millimeters long. (highly He5 332.8 1.9 1.7 In actual pellets, however, the length did not naturally 2530 drawn) exceed that of the pellet (about 2 mm). The two blends depicted in Fig. 3 were used in all * D.R. = drawratio = &/Ss further processing steps. The aim of the further studies E = elastic modulus; U = maximal strength; e = elongation; d = diameter of the strand. was to investigate the morphologies after extrusion and

Based Composites

627

I . Extrusion The morphology of the highly drawn blend after extrusion at different temperatures is depicted in Fig. 4. The effect of processing temperature on the morphology was similar to that found inthe preliminary tests. At the highest temperature (280°C Fig. 4e) both blend components were molten during processing aand fibrous morphology with fine well-oriented fibrils was formed in situ during the rapid coolingof the ordered melt. Elongationaldrawing of the solidifying extrudate increases the orientation of the LCP phases. Additional drawingat temperatures between the glass transition (TB) and meltingtemperature (T,) of the LCP phase, i.e., cold drawing, mayfurther improve the fibrillation and orientation. At moderate temperatures (230-250°C; Figs. 4c,d) the LCP fibers

LOW DRAW RATIO (L), 290 "C, (a)

CORE

were softened and agglomerated to clusters. At low temperatures (180-200°C; Figs. 4a,b)the fiber structure created during the previous melt blending step was fairly well maintained, as expected, but fiber orientation was poorer than in the blend processed at 280°C. Additional drawing could not be applied to the composites generated at the lower temperature owing to the poor melt strength of the material. Thus, the enhanced fiber orientationandreinforcing effect achievedwith the melt blends could not be achieved with the composites. 2. InjectionMolding The morphology ofthe injection molded blends is shown in the optical micrographs of Fig. 5. After injection moldingat 180"C,just asbefore, the

HIGH DRAW RATIO (H), 290 "C, CORE (b)

, c

Figure 3 Twin-screw extruded PP-LCP blend processed at a melt temperature of 290°C with low- (left) and high-draw ratio (right). Upper micrographs are taken from the core and lower ones fromskin the region.

628

Heino and Vainio

H e l , 180 "C,

He2,200

(4

(b)

"C,

He4. 250 "C (d )

Figure 4 Optical micrographs from the skin region of the single-screw extruded strands processed at cylinder temperatures of (a) 18O"C, (b) 200"C, (c) 230°C (d) 250"C, and (e) 280°C.

629

Composites Based on Premixed PP/LCP Blends

Lil, 18

Lil. 180 "C.

(a)

(b)

Hil, 180 "C, SKI (a

Hi2,280 "C, SKI (0 Figure 5 Opticalmicrographs of specimens injection molded at core (left) and skin region (right). (Sample codes as in Table 1.)

180°C (a-d)and 280°C (e,f). Samples were taken from

630

blendpreparedwith the higher draw ratio exhibited clearly thinner, longer, and more oriented LCP fibers than the less drawn blend (Fig. 5a,c). Thus, the difference in fiber length was maintained when the material was processed at a temperature clearly below the T, of the LCP. From the highly drawn blend a composite structure with long highlyoriented fibers was achieved. At the higher temperature (280°C) a distinct skid core morphology was formed with thin fibers in the skin region and dispersed LCP particles in the core (Figs. 5e,f). At this temperature both the blend components were molten during injection molding the andpreviously created fiber structure was destroyed. However, in the melt stagethe LCP phases were stillat least partly elongated owingto thelong relaxation time of the LCP. The fiber morphology was recreated during cooling in the mold, taking place as in situ fibrillation in the skin region due to the high-shear forces near the mold wall. Thus, the highly ordered morphology created during the injection stage could be locked by rapid coolingin the mold (low mold temperature). It should be noted that very high mold temperatures could cause relaxation of the oriented LCP phases. The composites injection moldedat the lower temperature (180°C) didnot exhibit any skidcore effect, but rather contained fibers throughout.

B. Mechanical Properties

Heino and Vainlo

2. Injection MoldedSpecimens The mechanical properties of the injectionmolded blends and composites are shown in Table 3. The differences in fiber length andorientation in the composites prepared at 180°C (Fig. 5a-d) are reflected in the strength and modulus.The composite based on the highly drawn blend showed better strength and modulus than the composite based on the less drawn one. After injection molding, both melt blends (Fig. 5e,f) exhibited a higher level of modulus than the composites, but the effect of the preceding drawing diminished. Nevertheless, the elastic modulus of the highly drawn melt blend was slightly better owing to the more elongated LCP phases. By contrast, the value of the impact strength was almost twiceas large for the composites as for the melt blends. This significant difference is explained by the totally fibrillarstructure of the composites. Although the thin fibrilsin the skin regionof the melt blends were able to carry most ofthe load appliedin the flexural test, they could not resist the impact loadto the same extent. In both the blends and composites, the addition of LCP reinforced the PP matrix considerably. On the basis of the fibrillar morphology throughout the specimens, even better mechanical properties were expected for the composites than for the blends. The poorer than expected reinforcement was primarily due to the lack of adhesion between fiber and matrix.

C. Rheology 1. ExtrudedStrands Shear viscosities of the twin-screw blended materials The tensile properties of the extruded blends and comwere measured at 190°C and 290°C (Fig. 6), the same posites are presented in Table 2. Compared to the neat temperatures as the melt temperatures during processPP, a clear reinforcement was achieved after twin-screw ing: 190°C for the composites and 290°C for the melt blending. The reinforcingeffectwas even moreproblends. nounced with the higher take-up speed (H), evidently The rheological behavior of the blends and composdue to the extremely fibrillar morphology, as seen in ites was totally different. Addition of LCP reduced the Fig. 3. The changes inthe morphology of the highly drawn blends after single-screw extrusion (Fig.4)werereTable 3 Flexural and Impact Properties of Injection flected in the mechanical properties, which also varied Molded Specimens with the processing temperature regions. A good reinforcing effect was achieved with the composites proCharpy cessed well below the T, of the LCP (He1 and He2, in impact Table 2) and with the melt blend (He5). However, the E U EY strength composites processed at moderate temperatures (He3 (MPa) (MPa) Material (mm) (kJ/m2)a and He4) showed poor mechanical properties due to the 12 (1) not broken 39 (1) 1390 (50) PP (180°C) clusterlike morphology. 40 (1) 12 (1) 41 (5) (180°C) 1580 (40) Lil The melt blends exhibited the best mechanical prop12 (1) 39 (5) 45 (1) Hi1 (180°C) 1770 (40) erties, which could be still further improved with addi11 (1) 22 (3) 41 (1) Li2 (280OC) 1900 (150) .tional drawing. The composites Hel-He4 could not be Hi2 (280°C) 1970 (110) 42 (1) 10 (1) 22 (1) drawn to improve the mechanical properties. In the case of the melt blends, even higher draw ratios than used in a Dimensions of the unnotched test bar = 4 X 10 X 112 mm; span this study will increase the fibrillation and orientation = 70 mm. Standard deviations in parentheses. of the LCP phase leading to significant improvements E = elastic modulus; U = maximal strength; ey = elongation at yield. in strength and modulus [21,30]. ~~

Composites Based on Premixed PP/LCP Blends

631

9 1000:

100:

LOW draw ratio 190°C High drawratio 190°C -A- Low drawretio 290°C -A- High drawratio 290°C -0PP 190°C 10: -0- PP 29OOC -m-0-

1000

100

i 10

'

a . . .

'1

.

' '

"*"I

W '

'

" " " I

I

Shear rate(S-') Figure 6 Shearviscosities of thepremixedPP-LCP blends measured at 190°C and 290°C.

viscosity of PP in the melt blends, as reported earlier [24,44], but increased it inthe composites. Nevertheless, the increase in viscosity caused by the solid LCP fibers was surprisingly small and the material was still easily processable at the temperature normally usedfor PP. A probable explanation is the small dimensions of the LCP fibers relative to conventional glass fibers for example. The differences in fiber length induced by drawing did not affect the viscosity markedly. V.

CONCLUSIONS

Blends of polypropylene(PP) and liquid crystalline polymer(LCP) processed withoutmelting the LCP were compared with conventional melt processed blends. In a first stage, PP was blended with 20 wt% of LCP in a twin-screw extruder with the take-up speed varied to achieve blends with differentLCP fiber dimensions. In the second stage, these blends were processed both below and above the T, of the LCP by extrusion and injection molding. At lower temperatures (180-200°C) the material was processed without melting the LCP anda real composite structure with solid LCP fibers in the PP matrix was formed. When processing was done above the T, of the LCP (280"C), allthe material was molten during processing and a compositelike blend morphology wascreated in situ during cooling of the oriented melt phase. The composite and blend morphologies were nevertheless different inthe solid state. In the composite there

were LCP fibers throughout the sample, whereas the blends showed a clear skidcore morphology with long oriented LCPfibers in the skin region and dispersed LCP domains in the core. The differences in fiber length induced by different drawingspeeds in the blending stage were maintained inthe composites, but vanished during the melt processing at higher temperatures. At processing temperatures 230-250°C the LCP fibers did not melt but were softened and agglomerated and an undesirable morphology with LCP clusters was obtained. The twin-screw mixed blend with the higher draw ratio showed significant reinforcementcompared to the neat PP andto the less drawn blend, andthis was promising for the additional experiments. The melt blends exhibited the highest level of modulus, but the effect of the preceding drawing seen in composites disappeared after injection molding. In comparison to the melt blends the compositesexhibitedsignificantlyhigherimpact strength. The addition of LCP sharplydecreased the viscosity of PP in the melt blends, but increased it inthe composites. The increase in viscosity effected by the solid LCP fibers was nevertheless surprisingly small. Thisworkledus to conclude that preblended PP-LCP blends can be processed without melting the LCP fibers. The resulting composite showed interesting morphology. In particular, the impact strength, which often is poor for conventional uncompatibilized LCP melt blends, was significantly higher for the composites. Other mechanical properties of the composites were not on a satisfactory level perhaps due to the lack of adhesion betweenfiber and matrix.By the addition of appropriate coupling agents the adhesion could be improved. On the other hand, the length of the pregenetrated fibers might be increased simply by increasing the length of the pellets. Themainbenefitsachievedbymeltprocessing PP-LCP blends are decreased viscosity and increased strength and modulus, which can be further improved by elongational drawing. Owing to the in situ fibrillation of melt blends, the morphology ultimately depends on the final processingconditions. The processing method leading to composites is advantageous in achieving good impact strength for moldings. Nevertheless, as typical for all composites, the poor melt strength of the PP-LCP composites does not allow high-speed drawing inextrusion. Althoughthe added LCPdoes not act asa processing aidas it does in the melt blends, it does not dramatically increase the viscosity either. REFERENCES 1.

W.J. Jackson, Jr. and H. F. Kuhfuss, J . Polym. Sci., Polym. Chem. Ed., 14: 2043 (1976). 2. T A . Chung, Polym. Eng. Sci., 26: 901(1986). 3. K. F. Wissbrun, J . Rheol., 25: 619(1981).

Heino and Vainio 4. T. Schacht,Doctor's thesis, Rheinisch-Westfalischen Technischen Hochschule, Aachen, West Germany (1986). 5. Y. Ide and Z. Ophir, Polym. Eng. Sci., 23: 261 (1983). 6. S. Kenig, Polym. Eng. Sci., 29: 1136 (1989). 7. Z. Ophir and Y. Ide, Polym. Eng. Sci., 23: 792 (1983). 8. L. C. Sawyer and M. Jaffe, J. Mat. Sci.,21: 1897 (1986). 9. F. N. Cogswell, B. P. Griffin and J. B. Rose, US.Patent 4,386,174, to ImperialChemical Industries (ICI) PLC, UK May 31, 1983. 10. E. G. Joseph, G. L. Wilkes and D. G. Baird, Polymeric Liquid Crystals, (A. Blumstein, ed.), Plenum Press, New York & London, p. 197 (1985). 11. A. Siegmann, A. Dagan and S . Kenig, Polymer, 26: 1325 (1985). 12. G. Kiss, Polym. Eng. Sci., 27: 410(1987). 13. Isayev, A. I. andModic,M., Polym. Compos., 8: 158 (1987). 14. Weiss, R. A., Huh, W., and Nicolais, L., Polym. Eng. Sci., 2 7 684 (1987). 15. Apicella, A., Iannelli, P., Nicodemo, L., Nicolais, L., Roviello, A., and Sirigu, A., Polym. Eng. Sci., 26: 600 (1986). 16. Brostow, W., Dziemianowicz, T. S., Romanski, J., and Werber, W., Polym. Eng. Sci., 28: 785 (1988). 17. Carfagna, C., Amendola, E., and Nobile M. R.,International Encyclopedia of Composites, Vol. 2 (S. M. Lee, ed.), VCH Publishers, New York, p. 350 (1990). 18. Dutta, D., Fruitwala, H., Kohli, A., and Weiss, R. A., Polym. Eng. Sci., 30: l005 (1990). 19. Seppala, J. V., Heino, M. T., and Kapanen, C., J. Appl. Polym. Sci., 4 4 : 1051 (1992). 20. Heino, M. T. and Seppala, J. V., Int. J. Mater. Product Technol. 7 56 (1992). 21. Heino, M. T. and Seppala, J. V., Polym. Bull., 30: 353 (1993). 22. Blizard, K. G. andBaird, D. G., Polym. News, 12: 44 (1986). 23. Blizard K. G. and Baird, D. G., Polym. Eng. Sci., 27: 653 (1987). 24. Heino, M. T. and Seppala, J. V., J. Appl. Polym. Sci., 48: 1677 (1993). .-~ -, 25. Datta, A., Chen, H. H., and Baird, D. G., Polymer, 34: 759 (1993).

26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48.

Taylor, G. I., Proc. Roy. Soc., A138: 41 (1932). Taylor, G. I., Proc. Roy. Soc., A146: 501 (1934). Wu, S., Polym. Eng. Sci., 27: 335 (1987). Plochocki, A. P., Dagli, S. S . , andAndrews, R.D., Polym. Eng. Sci., 30: 741 (1990). Heino, M. T. and Seppala, J. V., J. Appl. Polym. Sci., 44: 2185 (1992). Amano, M. and Nakagawa, K., Polymer, 28: 263 (1987). Ramanathan, R., BlizardK., and Baird, D., SPEANTEC, 34: 1123 (1988). Nobile, M.R., Amendola, E., Nicolais, L., Acierno, D., and Carfagna, C., Polym. Eng. Sci., 29: 244 (1989). Malik, T. M., Carreau, P. J., and Chapleau, N., Polym. Eng. Sci., 29: 600 (1989). Blizard, K. G., Federici, C., Federico, O., and Chapoy, L. L., Polym. Eng. Sci., 30: 1442 (1990). Carfagna, C., Amendola, E., Nicolais, L., Acierno, D., Francescangeli, O., Yang,B., and Rustichelli, F., J. Appl. Polym. Sci., 43:839 (1991). Crevecoeur, G. and Groeninckx, G., Polym. Eng. Sci., 33: 937 (1993). Done, D., Sukhadia, A., Datta, A., and Baird, D. G., SPE ANTEC, 36: 1857 (1990). Beery, D.,Kenig. S.. and Siegmann, A., Polym. Eng. Sei.,-31: 451 (1991). . Brinkmann, T., Hock,P.,andMichaeli,W., SPE ANTEC, 3 7 988 (1991). Carfagna, C.,Amendola, E., and Nicolais, L., Int. J. Mater. Product Technol., 7: 205 (1992). Incarnato, L., Nobile, M., and Acierno, D., Makromol. Chem., Macromol. Symp., 68:277 (1993). Kato, S., Naito, T., and Nakamura, K.,Japanese Patent 05305626A2, toToppan Printing Co., Ltd., Japan November 19 (1993). Heino, M., Hietaoja, P., Vainio, T., and Seppala, J., J . Appl. Polym. Sci.,51: 259 (1994). Vainio, T. P. and Sevvala, J. V., Polym. Polym. Compos., l : 427 (1993). Kijava, J., Rundqvist, T., Holsti-Miettinen, R., Heino, M., and Vainio,T., J. Appl. Polym. Sci., 55: 1069 (1995). Holsti-Miettinen,R., Heino, M., and Seppda, J.,J . Appl. Polym. Sci., 57: 573 (1995). OVectra Flussigkristalline Polymere(LCP),brochure from Hoechst Celanese, September 198911. Auflage. "

42 Compatibilization of Thermoplastic Elastomer Blends R. Asaletha* Rubber Research Instituteof India, Kerala, India

Zacharia Oommen and Sabu Thomas Mahatma Gandhi University, Kerala, India

1.

INTRODUCTION

During the last several years researchers all over the world have been trying to produce new polymeric materials with specific properties for specific applications. Since the field of new materials is highlyexposed, scientists have turnedtheir interest to modified forms such as polymer blends, polymer composites, interpenetrating networks, etc. As a new and important challenge for researchers, polymer blends have gained muchinterest and, of course, have becomea new branch of macromolecular science. The blending technique is quite attractive due to the fact that already existing polymers can be used and, thus, the costly development of new polymersviacopolymerization or polymerization of new monomers is avoided. Since there is no generally accepted definition of polymer blends, they are generally considered as a combination of two or more polymers. Polymer blendscan be obtained by mixingthe two polymers in the molten state, casting from common solvent, or latex blending,etc., and these methods do not involve the formation of chemical bonds. Blending is the simplest and cheapest process of combining the properties of different polymeric materials. We can classify blendsinto three categories: miscible, partially miscible, and immiscible. Miscibility can be defined in thermodynamicterms. For a binary blend to be miscible the following two conditions should be satisfied: AG,,, < 0 (1)

* Current aflliution: Mahatma Gandhi University, Kerala, India. 633

where AG, is the free enthalpy of mixing per unit volumeand is the volume fraction of component 2. In miscible polymer blends, molecular level mixing of the components is obtained. They are characterized by single-phase morphology. Immiscible blends do not satisfy either of the above criteria and will show a two-phase morphology. In the case of partially miscibleblends, the second criteria is not satisfied andwill show either twophase or single-phase morphology.

II. INCOMPATIBILITY

A. Problems and Solutions Only a very few polymers form truly miscible blends. These include poly(pheny1ene oxide)(PPOPS), poly(vinyl chloride) (PVC)-polymericplasticizers [l]. The rest of the blends, which are either partially miscible or immiscible, may undergo micro- or macrophase separation, leadingto heterophase polymer blends.This heterogeneity is an unfavorable one, and this often leads to problems that reflect in the overall performance of the resultant material. Blendingcan give rise to morphologies that lead to certain specific characteristics. It is expected that this process can give rise to a material with the proper balanceof properties than would be obtainable with single polymers.Practically, it is difficultto get the expected combination of properties due to the fact that many of the polymers are thermodynamically immiscible and it is difficult to get a homogeneous product. In an immiscible blend, the situation at the blend interface is critical, i.e., a high-interfacial tension and poor adhesion between the phases are observed. The high

634

Asaletha et al.

0

0

side by side

Figure 1 Various phase configurations uponthe extrusion of binary immiscible blend.

viscosities associated with such systems are responsible for poor dispersion and lack of stability to gross phase segregation. The various phase configurations of a binary blendas suggested by Paulet al. [2] are given in Fig. 1. The low-intermolecularforce between the component phases, which is responsible for the poor properties of incompatible blends, can be improvedby increasing the interfacial area and adhesion between the phases. The term “compatibility” is usedextensively in the blend literature and is used synonymously with the term “miscibility” in a thermodynamic sense. Compatible polymers are “polymermixtures that do not exhibit gross symptoms of phase separation when blended” or “polymer mixturesthat have desirable chemical properties when blended.’’However, in a technological sense, the former is usedto characterize the ease of fabrication or the properties of the two polymersin the blend [3-53. 111.

THERMOPLASTICELASTOMER BLENDS (TPEs)

TPEs are materials that possess, at normal temperatures, the characteristic resilience andrecovery from the extension of crosslinked elastomers and exhibit plastic flow at elevated temperatures. They can befabricated by the usual techniques such as blow molding, extrusion, injection molding,etc. This effect isassociated with certain interchain secondary valence forces of attraction, which have the effect of typical conventional covalent crosslinks, but at elevated temperatures, the secondary

bonds dissociates and the polymer exhibits thermoplastic behavior. TPEs have many processing advantages over the vulcanized rubbers. The complexvulcanization techniques are not required and very little compounding is needed. Thermoplastic processing techniques, such as blow molding,heat welding, etc., that are unsuitable for conventional rubbers canbeappliedsuccessfully to TPEs. The short processing cycle and mixing will result in very low-energy consumption. Processing of scraps that are considered as a waste in a conventional system can be recycledhere. Besides allthese advantages, certain disadvantages were also observed, i.e., they show high creep and set on extended use and will melt only at elevated temperatures. Thermoplastic elastomers have becomevery important. The estimated use of TPEs world widein 1985 was 500,000 metric tonnes. They are replacing many of the conventional rubbers as well as thermoplastics. TPEs canbemainlyclassified into five groul These include: 1. TPEs from rubber-plastic blends. 2. Polystyrene-elastomerblock copolymers. 3. Polyurethane-elastomerblock copolymers. 4. Polyamide-elastomerblock copolymers. 5. Polyether-elastomerblock copolymers.

TPEs from blends ofrubber and plastics constittl1 an important category of TPEs. These can be prepared either by the melt mixing of plastics and rubbers in an internal mixeror by solvent casting froma suitable solvent. The commonly used plastics and rubbers include polypropylene (PP), polyethylene (PE), polystyrene (PS), nylon, ethylene propylene diene monomer rubber (EPDM), naturalrubber (NR), butyl rubber, nitrile rubber, etc. TPEs from blendsof rubbers and plastics have certain typical advantages over the other TPEs. In this case, the required properties can easily be achievedby the proper selection of rubbers and plastics and by the proper change in their ratios. The overall performance of the resultant TPEs can be improved by changingthe phase structure and crystallinity of plastics and also by the proper incorporationof suitable fillers, crosslinkers, and interfacial agents.

IV. COMPATlBlLlZATlON CONCEPTS Although blendingis an easy method for the preparation of TPEs, most of the TPE blends are immiscible. Very often the resulting materials exhibit poor mechanic properties due to the poor adhesion between the phases. Over the years different techniques have been developed to alleviate this problem. One way toisalter the blending technique so that the interfacial area between the component phases can beincreased. By the proper selection of the ’processing technique either a co-continuous or

Compatibilizationof Thermoplastic Blends

635

Table 1 CompatibilityThroughNonreactive

Table 3 Compatibility Through Low-Molecular Weight

Copolymers

Minor

Reactive Compounds

Major Compatibilizer component component PE or PSPS PP PE or PP EPDM PS

HDPE

PET

Minor

Major Compatibilizer component component

S-B, S-EP,S-I-S, S-I-HBD, S-EB-S, S-B-S, PS-PE graft copolymers PS or PMMA S-EB-S EPM, EPDM PP or PE PMMA EPDM-g-MMA PA-6 or EPDM PS-PA-6 block copolymers or PPE or S-EB-S S-EB-S or PE

Fluororubber, NBR or CHR Triazine dithiol complex FPM PVC or LDPE or PVC Polyfunctional monomers plus peroxide LDPE NBR curative and NBR PP interchain copolymer Cholrinated paraffh PP or PVC PVC or PP Amino silane PPE PA-6,6 Source: Ref. 1.

Source: Ref. 1.

interpenetratingphase morphologycan be obtained that results in direct load sharing withoutthe need for stress transfer across the phases. The second way is by the addition of a third componentthat is ableto have interaction with blend components, (e.g., block and graft copolymersandlow-molecularweightmaterials). The third technique isto blend suitably functionalized polymers capable of specific interaction or chemical reactions. This functionalization can be done in solution or in a compoundingextruder [6] and may involve reactions such as halogenation, sulfonation, hydroperoxide formation, and the in situ formation of block and graft copolymers. The “in situ” formed copolymers act as very good compatibilizers in many systems. These are formed during compounding, mastication, polymerization of one monomer in the presence of another polymer, etc., and have segmentsthat are chemically identicalto the homopolymers. Hajianet al. [7] reported the in situ formation of styrene-ethylene graft copolymers duringthe mixing of PS and PE. Anderson [8] studied the compatibilizing action of in situ-formed EPDM-g-MMA duringthe melt

Table 2 Compatibility Through Reactive Copolymers

PP or PS-6 PE

PA-6-PA-6,6 copolymer PA-6 or PP PA-6 or PA-6,6

PP or PE

PET

Source: Ref. 1.

V. COMPATIBILIZERS

A. Requirements

Major Compatibilizer component component ABS

extrusion of EPDM and methylmethacrylate (MMA). The in situ formationof compatibilizers usingfunctionalized polymers formthe subject of several studies. Ide and Hasegawa[9] studied the effect of maleicanhydride grated polypropylene in PP-polyamide-6 blend. Block or graft copolymers that act as compatibilizers are of two types, reactive and nonreactive. Nonreactive ones have segments capable of specific interaction with each of the blend components. In reactive copolymers, segmentsare capable of forming stronger covalent or ionic bonds withthe blend components. Copolymers of both A-B type and A-C type can act as efficient compatibilizers in the A-B system, provided C is miscible with B. Tables 1 and 2 contain a few examples of polymer systems that are compatibilized through reactive and nonreactive and nonreactive copolymers, respectively. Systems withlow-molecularweightcompounds as compatibilizershavebeen subjected to various studies in the field of blends. Co-crosslinking, crosslinking, and grafting reaction may be involved in such systems and may lead to the formation of certain copolymers. Table3 deals with such a system involving low-molecular weightcompounds as compatibilizers.

SAN-MA copolymer EPM-MA copolymer Ionomers, carboxyl functional PES PP-g-AA, Carboxyl functional PE

The efficiency of the copolymers, either block or graft, acting as the compatibilizer depends on the structure of the copolymers. Oneof the primary requirements to get maximum efficiencyis that the copolymer should be located, preferentially at the blend interface (Figs. 2a, b, and c). There are three possible conformations, as shown in the figure. Many researchers [lo-121 found that the actual conformation is neither fully extended nor flat (Fig. 2c). A portion ofthe copolymerspenetrates into the corresponding homopolymer and the rest re-

636

Asaletha et al.

Figure 2 Conformation of copolymer at the blend interface: (a) completely extended, (b) completely flat, and (c) neither completely extended nor completely flat.

Compatibilization of Thermoplastic Blends

637

mains at the interface. In thisrespect, a block copolymer is supposedto be superior to a graft copolymerbecause a less-complicated conformation expected is in the former case [13,14]. In the case of graft copolymer, multiple branches should be avoided because they would restrict the penetration of the backbone into the homopolymer phases, and a diblock copolymer will be more effective than a triblock copolymer.Fayt et al. [l51 demonstrated that a tapered diblock isore efficient than a pure diblock with the same composition and molecular weight.Pure diblockcopolymer contains highlyincompatible sequences. These sequences segregate into domains and no mixingoccurs. But tapered block copolymers do not formdomains of their ownand therefore provide a strong adhesion. Compared with diblock, the tapered block copolymers can be easily dispersed due to their low viscosity. Chemical identity of the copolymer segment with the homopolymer phase is important. Even though there is no chemical identity between the copolymer segment andthe homopolymer, copolymers can be kept equally efficient provided the segment is miscible with the homopolymer. Another important requirement is that the copolymer should have the propensity to segregate into two phases. In addition, the copolymer, both block andgraft, should not be miscible as a whole in one of the homopolymer phases. The amount of the copolymer to be added into a binary blend depends on several factors. Since the copolymer preparation is an expensive process, it should be used to its maximum efficiency. The amount of copolymer (m)required to saturate a unit volumeof blend is given by: m

=

3tp~MlaRN

(3)

where t p is ~ the volume fraction of polymer A , R is the radius of dispersed particle A in a matrix B, N is the Avogadro’s number, and M is the molecular weight of the copolymer. For a copolymer to be fully efficient, its molecular weight (M) should be higher than the molecular weight of the homopolymers. Riess and Jolivet [l41 studied the effect of molecular weight on solubilization. When the homopolymer molecular weight is larger than that of the corresponding blocksegment, the homopolymer forms a separate phase and is not solubilized into the domains of the block copolymer. Inthe case of highmolecularweight copolymers, the long segments are able to anchor the immiscible phases firmly. Besides all these arguments, there should bean optimum molecular weight when the cost-benefit of the resultant material is a concern. The effect of the molecular weight of poly(styrene-butadiene) copolymer in PS-PB system has been explained by Legge et al. [16]. They have shown the interfacial profiles of two systems: (1) with block copolymer having molecular weight less than that of the homopolymer, (2) with blockcopolymer having molecular weight higherthan that of the homopolymer. It was

found that the copolymer with high-molecular weight has greater concentration in the interfacial region. The emulsifying effectof block and graft copolymers in polystyrene-polyisoprene (PS-PI) blends was demonstrated by Riess and Joliviet[14]. Mechanical properties of the incompatible blendscan be enhanced by controlling the dispersed phase size and adhesion between the components, which can be achieved bythe addition of a suitable copolymer. Molecular weight and composition are two important parameters that determine whether the copolymer will locate at the blend interface, in the continuous phase, or in the dispersed phase. Block copolymers of equal segmental mass are more effective as a compatibilizer than those of unequal segmental mass. There are various techniques for locating the copolymer. These include the use of a copolymer witha fluorescent group,x-rayscanningmicroanalysis, analysis of gel formed after crosslinking the elastomer phase by yradiation, etc. In additionto the abovementioned parameters, various factors such as viscosity of the copolymer and its interaction with the homopolymers also play a major role in the compatibilization process. VI. THEORETICALASPECTS OF COMPATlBlLlZATlON Leibler [l71 and Noolandi et al. [18,191 developed thermodynamic theories concerning the emulsificationof copolymers (A-b-B)in immiscible polymer blends (A-B). The theory of Leibler holds for mainly compatible systems. Leibler developed a mean field formalism to study the interfacial properties of two polymers, A and B with an A-B copolymer. An expression for interfacial tension reduction was developed by Noolandi and Hong [l81 based on thermodynamics to explain the emulsifying effect of the A-b-B in immiscible A-B blends (A-A-bB-B). [18,19]. The expression for interfacial tension reduction ( Ar ) in a binary lend uponthe addition of divalent copolymer is given by:

Ar = dtp,(l/zx

+ 1IZc) - 1IZc exp( Zcx/2)

(4)

where d is the width at half height of the copolymer reduced by Kuhn’s statistical segment length, 4, is the bulk volume fraction of the copolymer in the system, x is the Flory-Huggin’s interaction parameter between A and B segment of the copolymer, and Zc is the degree of polymerization of the copolymers. Althoughthis expression was derived for the action of symmetrical diblock copolymers (A-b-B)inhomopolymerblend (A-B), i.e., A-A-b-B-B, this theory can be successfully applied to A-A-g-B-B system. According to this theory, the interfacial tension reduction shoulddecrease linearly withcopolymer content at low concentrations followed bya leveling offat higher concentrations. The theory of Noolandi and Hong [ 181

638

Asaletha et al. I

has been testifiedto by Koberstein and coworkers [19a] and Thomas[101. COMPATlBlLlZATlON OF TPEs

VII.

A. Specific Cases There are a large numberof studies related to the compatibilization of TPEs by the addition of copolymers (both block and graft) and by dynamic crosslinking. Park et al. [20] reported on the synthesis of poly(chloroprene-co-isobutylmethacrylate) and its compatibilizing effect in immiscible polychloroprene-poly(is0butyl methacrylate)blends, A copolymer of chloroprene rubber (CR) and isobutyl methacrylate (iBMA) poly[CPCO-(BMA)] and a graft copolymer of iBMA and polychloroprene [poly(CR-g-iBMA)] were prepared for comparison. Blends of CR and PiBMA are prepared by the solution castingtechnique using THF asthe solvent. The morphologyand glass-transition temperature behavior indicated that the blend is an immiscible one. It was found that both the copolymers can improve the miscibility, but the efficiency is higher in poly(CR-COiBMA) than in poly(CR-g-iBMA). Oommen andThomas [21] studied the interfacial activity of natural rubber-g-poly(methylmethacry1ate)in incompatible NR-PMMA blends. Graft copolymer of NR and PMMA was prepared using a redox initiator consisting of cumene hydroperoxide and tetraethylene pentamine. Mechanical and morphological analyses of the blendswithandwithout the compatibilizer were studied, and it was found that the mechanical properties increase with increasing loadingof the graft copolymer concentration (Table 4). Morphologicalanalyses are in agreementwith the mechanical data. Figure 3 gives the scanning electron micrograph(SEM) of 60:40 NP-PMMA blendscontainingdifferentlevels of the compatibilizer. Theyfurther studied the effect of casting solvent, mode of addition of compatibilizer, molecular weight of homo- and copolymers, etc., on the morphological and mechanicalproperties. The experimental results were compared with the theoretical predictions of Noolandi and Hong[ 181.

I

(b) Figure 3 Scanning electron micrographs

of 60:40 NR-PMMA blend: (a) 0% and (b) 3.5% graft copolymer.

Asaletha and coworkers [12,22] further studied the compatibilizingeffect of NR-g-PSinNR-PS blends. NR-PSblend is incompatible and immiscible and its compatibility can be improved by the addition of the

Table 4 Effect of Compatibilizer Loading on Mechanical Properties

Percent of at Stress 50% strength graft elongationelongationelongation copolymer (N/mm) 0

1

2.5 3 Source: Ref. 21.

362 429 262 381

1.os 1.07 2.55 1.18

Stress at 100% (N/mm)

1.41 3.98 1S4 6.97 7.433.92 7.811.95

at Stress Elongation Tensile 130% (N/mm)

1.70 1.96 4.55 2.56

(N/mm2)

( m

Compatiblllzation Blends of Thermoplastic

639

Table 5 Mechanical Properties of 50:50 NR-PS Blends ~~

~~

Wt% of graft polymer 1

0 1.5 3 4.5 6 7.5

Stress at 15% elongation

(MW

2.37 1.24 2.45 S 4 2.75 3.07 3.24 3.88

Stress at 30% elongation W")

Stress at 50% elongation

Tensile strength

( m

( M(WM W

1.78 1.82 2.05 2.28 2.78 3.47

Elongation at break

3.60 3.86 4.50 10.10 13.24 13.15

0.30454 194 2.10 190 1.63 252 1.39 247 1.37 1 24

Tensile impact strength (J/m2) x 1.64 x x x x

lo' lo' lo'

lo' lo' X lo'

Source: Ref. 12.

graft copolymer. Up to a particular concentration of the compatibilizer, the mechanical properties of the blend

increased, thereafter the properties leveled off at higher graft loading (Table 5). The compatibilizing effect de-

pends upon the solvent used for the preparation of the blend (Fig. 4), which turn, in depends upon the solubility parameter of the concerned solvent. The molecular of prepaweight of copolymer and homopolymers, mode

Figure 4 Effect of casting solvent on compatibilization studies of NR-PS blend. (a) 0% graft (CHCls), (b) 1.2% graft (CHCls), (c)0% sraft (CCL), and (d) 1.2% graft (CCL).

Asaletha et al.

640

ration of the blend, composition of the blend, etc. are other parameters that control the compatibilization process. Morphology and mechanicalproperties were studied in detail andthe area occupied by the compatibilizer in each case was calculated and correlated with the theoretical predictions of Noolandi and Hong [18]. Characterizationand control of interfaces in the incompatible polymer blends were reported by Fayt et al. [23]. Theyused techniques such as electron microscopy, thermaltransition analysis, and nonradiative energy transfer (NRET), etc. They have illustrated the exciting potentialities offeredby diblock copolymers in high-performance polymer blends. Chu et al. [24] correlated viscosity-morphology and compatibility ofPS-PB blends. The effect of styrene-butadiene triblock copolymerin PS-PB was studied, and it was found that the domain size decreases with an increase of compatibilizer loading.The blending methods influenced the morphology due to the difference in the extent of mixing. Dynamic crosslinking as a means to improve the impact strength and other mechanical properties of polypropylene-elastomer blends has been discussed in detail by Inoue [25]. All these blends contain 80% PP and 20% elastomer. Elastomers includeEPDM, styrene-butadiene-styrene (SBS) and styrene-isoprene-styrene (SIS)and the crosslinking system comprised of N,N’-mphenylene-bis-maleimide and 6-ethoxy 2,2,4-trimethyl1,Zdihydroquinoline or poly(2,2,4-trimethyl-2,2-dihydroquinoline).Impact strength and other mechanical properties, such as tensile strength at yield,ultimate elongation, flexural modulus, etc., showed remarkable increase after crosslinking. This is due to the increase of interfacial adhesioncaused by the PP graft elastomers located at the blend interface. Interfacial adhesionand, thereby, compatibility can be enhanced by the selective crosslinking reaction in polymerblends. Inoue and Suzuki [26] reported the properties of blends dynamicallycrosslinked PP-EPDM blends. The crosslinking agent wasNN-m-phenylenebismaleimide - poly(2,2,4- trimethyl - 1,Zdihydroquino line) system. Increase in interfacial adhesion leads to

Table 6 Effect of Cure Inhibitor Concentration on the Charpy Notched Impact Strength(ad at 0°C

CTPI (min) (phr) 0 0.3 0.6 1.o

Emin t,, tm, - tdn (min)80120 85/15(min)

0.3 0.6

1.o

2.0

Source: Ref. 21.

1.7 2.9 3.3 4.3

1.4 2.3 2.3 2.3

ak, kJm-2

13.9 14.2 17.0 19.1

24.1 38.2 39.4 37.7

an improvementin Izod impact strength. Various other mechanical properties such as tensile strength at yield, ultimateelongation,andflexuralmodulus were also studied before and after the crosslinking reaction. Krulis et al. [27] also described the dynamic crosslinking as a route to improve the mechanical properties. It was found that a high-impact strength is obtained in PP-EPDM blends by slow curing with sulfur. Thiuram disulfide N-(cyclohexy1thio)phthalimide was used as an inhibitor of curing, and its effect on the impact strength ofdynamically cured PP-EPDMblendswas studied (Table 6).It was also found that the one-step method of blend preparation also has a favorable effect on the impact strength of the resultant blend system. Compatibilization along with dynamic vulcanization techniques have been used in thermoplastic elastomer blends of poly(buty1eneterephthalate) and ethylene propylene diene rubberby Moffett andDekkers [28]. In situ obtained by the use formation of graft copolymer can be of suitably functionalizedrubbers. By the usage of conventional vulcanizing agents for EPDM, the dynamic vulcanization of the blend can be achieved. The optimum effect of compatibilization along with dynamic vulcanization can be obtained only whenthe compatibilization is done before the rubber phase is dispersed. Ha [29] has shown that in PP-HDPE (of high-density PE) dynamically cured EPDM, the cured EPDM act as a compatibilizer to the HDPE-PP system. Blending was done in two ways. EPDM wascured first and then blended with PP andHDPE. In the second case, EPDM was cured in the presence of PP and HDPE using dicumy1 peroxide (DCP)as the vulcanizing agent. In EPDMrich composition, mechanicalproperties were increased by increasing the concentration of DCP, whereas in PPrich composition, the reverse was the case. Dynamic vulcanizationas a method to improve the mechanical properties of NR-PE blends has been discussed in detail by Choudhary et al. [301. The physical properties of unvulcanized and vulcanized NR-HDPE blends are given in Table 7 where notations A, B, and C indicate 70:30, 50:50, 30:70 NR-HDPE blends, respectively. The subscripts C and D denote blends containing DCP and highabrasion furnace (HAF) black (40 phr), respectively. Inall the compositions, the DCP-curedblends showed better properties than the corresponding unvulcanized samples. Choudhary et al. [301 further demonstrated the use of EPDM, chlorinated PE, chlorosulfonated PE, maleic anhydride modified polyethylene, and blends of epoxidized natural rubber-sulfonated EPDM as compatibilizers in NR-LDPE (low-density PE) blends. It was foundthat PP-EPDM blends [30] with a slow curing EPDM have a high-impact strength. Effect of tetramethyl thiuram disfulfide (TMTD) concentration on the Charpy notched impact strength of PP-EPDM blend is given in Table 8.

TMTD

Compatibilization Blends of Thermoplastic

641

Table 7 Physical Properties of NR-HDPE Blends

HDPE Property

A

13.4 29.4 (MPa) modulus Young's Modulus 100% (MPa) strength (MPa) Tensile 11.310.914.613.532.1 Elongation at break (%) 30.0 strength Tear 40.828.1 (kNm") 55.747.274.472.5 118.0 setTensile (%) -

Ac

65.0

-

1.4 14.5 6.1

-

B

Bc

5.8 7.2

8.7

24.0 960 430 420

-

92

>l00

cc

C

cb

NR

cbc

2.4 2.1 4.5 2.2 4.65.7 4.1 2.3 12.7 9.5 13.7 13.0 500 450 300 470 43.6 4.3536.4 72 64 70 68

0.4 0.8

>l00

Source: Ref. 30.

A binary blend of polypropylene-ethylene propylene dienemonomer rubber and a ternary blend of PP-EPDM-poly(ethy1ene-co-methacrylic acid) ionomer was prepared by Kim et al. [31]. The rheological, mechanical,andmorphological properties of these blends were analyzed. Two kinds of ionomers neutralized with different metal ions (Na+ and Zn2+)were used and their concentration varied from5 2 0 % based on the total amount of PP and EPDM. It was found that rheological and morphological properties of the binary and ternary blends showed much variation due to the compatibilizing effect of the ionomer. Na neutralized ionomer (ionomer A) showed a compatibilizing effect, B) did not, whereas the Zn neutralized ionomer (ionomer and the effect was prominent at 5 wt% of the ionomer concentration. The effect of ionomer concentration on the storage modulus of PP-EPDM blends is given in Fig. 5. It was seen that the ternary blends with 5 wt% of ionomer A showed better properties than other blends. The effect of ionomer concentration on the mechanical properties of PP-EPDM blends is given Table 9. It is seen that the tensile strength and modulus show a maximum at 5 wt% of both ionomerA and B,thereafter, it decreases at higher ionomer loading. The properties are higher for ternary blends containing ionomerB than these containing ionomer A. On the other hand, addition

of ionomer will reduce the elongation at break regardless of the ionomer type. The structure-property relationship of graft copolymers based on an elastomeric backbone poly(ethy1acrylate)-g-polystyrene was studiedby Peiffer and Rabeony [32]. The copolymer was prepared by the free radical polymerization technique and, it was foundthat the improvement in properties depends upon factors such as the number of graftdchain, graft molecular weight,etc. It was shownthat mutually graftedcopolymers produce a variety of compatibilized ternary component blends. Coran and Patel [33] selected a series of TPEs based on differentrubbers and thermoplastics. Three types of rubbers EPDM, ethylene vinyl acetate (EVA), and nitrile (NBR) were selected and the plastics include PP, PS, styrene acrylonitrile (SAN), and PA. It was shown that the ultimate mechanical properties such as stress at break, elongation, and the elastic recovery of these dynamically cured blends increased with the similarity of the rubber and plasticin respect to the critical surface tension for wetting and with the crystallinity of the plastic phase. Critical chain lengthof the rubber molecule, crystallinity of the hard phase (plastic), and the surface energy are a few of the parameters used in the analysis. Better results are obtained witha crystalline plastic material when the entanglement molecular length of the

Table 8 Effect of Composition of TMTD Curing System on the Charpy Notched Impact Strength (ak) for85/15 and 80/20 PP/EPDM Blends

ZnO 1.5 3 3 3 3 Source: Ref. 30.

(Phd

(PW

0.5 1 3

0.5 1 1 2 3

3

10.1

ak, Urn-*

(min)

tmin (min) 80/20

85/15

12.1 7.2 3.5 4.0 7.0

24.7 5.7 27.3 17.8 3.1 3.6 23.7 5.5 23.7

12.4 15.6 12.0 13.8 17.6

tmin

tmax

(min) 2.0 1.5 0.4 0.4 1.5

Lax

-

642

Asaletha et al.

r

l

10.

10'

A

10'

1 Q"

I 1

l0

. *..

10'

10. 20 -'

1

10

10'

Frequency (rawsec)

Figure 5 Effect of ionomer concentrationon storage modulus of PP-EPDM blends.(a) PPSO-EPSOhA, and(b) PPSO-EPSO/ IB (A)0 wt%, (0) 5 wt%, (0) 10 wt%, (A) 15 wt%, and (m) 20 wt%.

rubber is low. The other two conditions that are to be satisfied are: (1) the surface energies of the plastic and rubber shouldnot be very high, and(2) both rubber and plastic should not decompose in the presence of each other during melt mixing. In the case of NBR-nylon blends it was reported that the addition of the curative m-phenylene-bis-maleimide, improvedthe strength and stiffness of the blend. An attempt to compare the effect of different curative

Table 9 Effect of Ionomer Concentrationon the Tensile Propertiesof Blends of PP-EPDM ~~

Sample PPSO-EPSO PPSO-EPSOhAS PPSO-EPSOhAlO PP50-EPSOhS15 PP50-EP50/IA20 PPSO-EPSO/IBS PPSO-EPSO/IB 10 PPSO-EP50/IB1.5 PP50-EP50/IB20 Source: Ref. 31.

Modulus, E (Kgf/cmz) 575.82 630.21 497.35 487.49 421.72 719.63 680.21 640.81 611S8

Elongation Tensile at break strength (Kgf/cm*) (%l 105.57 109.23 88.08 74.68 56.23 112.62 93.07 85.54 65.50

234.75 138.60 170.02 237.00 190.43 195.18 206.33 215.45 176.66

m-phenylenebismaleimideanddimethyl01phenolic compounds wasalso reported. Nylon-NBR graft molecules formed during the crosslinking induce better homogenization in the system, which leads to overall enhancement of the blend performance. In additionto dynamic vulcanization,the technological compatibilization technique was also adopted by Coran and Patel [34] to obtain thermoplastic vulcanizate having good mechanical integrity and elastic recovery. Swelling of thermoplastic elastomeric vulcanizates using a model EPDM-PP blend in various solvents, such as cyclohexane, butyl acetate, methyl ethyl ketone (MEK), etc. was studied by Coran andPatel [35]. Blends were vulcanized both by dynamic and by static means and the mechanical properties, such as ultimate tensile strength, Young's modulus, ultimate elongation, hardness etc., were determined ineach case. All the properties except elongation break (EB) are higher for dynamically vulcanized samples. Static samples are not processable with the typical thermoplastic processing techniques, whereas the dynamicsamples can be molded. The amount of swelling of the thermoplastic vulcanizates was found to be less than the average swelling of the rubber and plastic. It was also noted that the vulcanizates prepared by the dynamic technique will swellless than the composition prepared by static means. Legge et al. [36] have discussed the PS-PI blend and the corresponding PS-b-PI as a model system for

Compatibilizationof Thermoplastic Blends

643

rubber-modified thermoplastics. The emulsifying effect of the block copolymer is evaluated by checking the transparency of the polymer blend.The transparency of an incompatible PS-PI system having different refractive indices can be obtained by reducingthe particle size of the dispersed phase below a certain level. This may be possible by the compatibilizing action of the block copolymer. Frounchi and Burford [37] studied the effect of styrene block copolymer as a compatibilizer in isotactic PP-ABS blends. It was found hat in PP-rich blends a marginal improvementin mechanical properties was obtained. However, in acrylo nitrile butadiene styrene (ABS) rich blends no improvement was obtained. The effects of four different block copolymers, SBS, SIS,

Mechanical Properties of PP-Rich Blends

Yield tensile Elongation at break, EB strength, uy (MW (%l

PP80-ABS20 PP80-ABS20-SEBSS PP80-ABS20-SEB5 PP80-ABS20-SIS5 PP80-ABS2O-FSEBS5 23

19

5

22 24 22 25

10 15 17

Mechanical Properties of ABS-Rich Blends Composition of blend (ABS-rich)

Source: Ref. 37.

ABSSO-ABS20 ABS80-PP20-SEBS5 ABS80-PP2O-SEB5 ABS80-PP20-SISS 23 ABS80-PP20-FSEBS5

19

5

22 24 22 25

10 15 17

Source: Ref. 37.

7.0

n n

Yield tensile Elongation strength, U, break, at EB (MW (%l

styrene-ethylene-butylene-styrene (SEBS), and a slightly maleated functionalized SEBS were compared. Mechanical properties of the PP-ABS blends with different block copolymers are given in Tables 10 and 11. The effect ofABSwt% on the Charpy impact strength is given in Fig. 6. Elliot [38]has reported that interfacial adhesion in the NR-PP blend can be enhanced by the addition of small amounts of HDPE. Addition of HDPE does give some improvementin the notched Izod impact strength of NR-PP blend (Fig. 7). The effect of HDPE on the impact modification of NR-PP is associated with the improved crystallinityof PP, enhanced by HDPE. During the mill mixing of NR and PP, chain scission may occur to give polymeric radicals that, on reaction with

Table 10 Effect of Dif'ferent Block Copolymers on the

Composition blend of (PP-rich)

Table 11 Effect of Different Block Copolymers on the

4

I '

.

u.u - 4

0

20

40 ABS

1

60

1

80

I

I

100

b ' t %)

Figure 6 Effect of ABS weight percent on Charpy impact strength.

644

Asaletha et al.

lzod impact strength (J m ')

300

lzod impact strength (J m")

-

200

100

-

-50

-30

-10

Temperature PC)

10

Figure 7 Effect of additionofHDPEonIzodimpact strength of NR-PP blend. (0) 20:80 NR-PP homopolymer, (V) 20:67: 13 NR-PP-homopolymer-HDPE, (U)15:85 NR-PP copolymer grade, and(x) 15:75 : 10 NR-PP-copolymer-HDPE.

the addedmultifunctionalradical

acceptor, cangive

graft copolymer and, in turn, can act as a compatibilizer in a NR-PP system. The effect of such compatibilizers

I

0-50

-io

-io -io - i o

b' tb

Temperature (*C)

Figure 8 Effect of addition of HDPE on BS-notched impactstrength ofNR-PPblend. (V) PP, (V) PP + 0.5 p.p.h.p. HVA2, (0)blend, (0) blend + 0.2 p.p.h.p. HVA2, (0)blend + 0.5p.p.h.p.HVA2,and(A)blend + 1p.p.h.p. HVA2.

on the BS-notched Izod impact strength is given in Fig. 8.

Compatibility and variousother properties such as morphology, crystallinebehavior, structure, mechanical properties of natural rubber-polyethylene blends were investigated by Qin et al. [39]. Polyethylene-b-polyisoprene acts as a successfulcompatibilizer here. Mechanical properties of the blends were improved upon the addition of the block copolymer (Table12). The copolymer locates at the interface, and, thus, reduces the interfacial tensionthat is reflectedin the mechanical properties. As the amount of graft copolymer increases, tensile strength and elongation at break increase and reach a leveling off. Morphological studies of these blends revealed that the compatibilization was very effective in decreasing the interfacial tension and increasing the adhesion between the two phases (Fig. 9).

Table 12 Mechanical Properties of NR-LLDPE

Compatibilized with PE-b-PI Tensile strength Elongation break at (kdm-')

Sample

240 84 99 112 132 113 110 96

0: 1oo:o

50:50:0

4950: 1 47:50:3 45:50:5 41:50:9 3750: 13 37:50:13"

" [v]PE-b-PI

=

Source: Ref. 39.

6 0 )

1387 975 965 1190

1089 1089 1188 935

12.8 dlg"; 18.9 dlg" for other blends.

Compatibilization Blends of Thermoplastic

645

L I

I

(B) Figure 9 (A) SEMmicrographs of NR-LLDPE-PE-b-PI : (a) 50:50:0, (b) 45:50:5, (c) 3750: 13, and (d) 37:50:13; (B) TEM micrographs of NR-LLDPE-PE-b-PI: (a) 50:50:0, (I3) 45:50:5, (c) 37:50:13, and (d) 37:50:13.

646

Asaletha et al.

The molecular weight of the copolymer (PE-b-PI) has a profound influence on the mechanical properties of these blends (NR-LLDPE) (Fig. 9). High-molecular weightcopolymeris a more effective compatibilizer compared with a low-molecular weightone. Improvement in the compatibility of NBR-PP blends by various compatibilizers was reported by Rader and Sabet [401. The effect of different compatibilizers on the properties of NBR-PP blends is given in Table 13. Wang and Chen[41] studied the compatibility problems of incompatibleNBR-PVC blends. Poly(viny1idene chloride-covinyl chloride) is reported to act asan efficient interfacial agent. Blends of PVC, NBR, andthe copolymer were prepared by the solution casting technique usingTHF asa solvent. Improvement in mechanical properties can be achieved in NBR-PVC blend by the addition of different types of rubbers [42]. Different rubbers include NR,styrene butadiene (SBR) and butadiene (BR). Replacement of a few percent of NBR by other rubbers will improve the mechanical properties and at the same time reduce the cost of the blend. Compatibility of immiscible PP-NBR blends was improvedby the reactive compatibilization technique using various modified polypropylenes. In this study,

Figure 10 Effect of compatibilizer molecular weight on mechanical properties. (A) NR-LLDPE-PE-b-PI = 50: 50:0, (B) NR-LLDPE-PE-b-PI = 37: 50: 13 (M, = 1.29 x lo5), and (C) NR-LLDPE-PE-b-PI = 37:50: 13 (M, = 2.79 x 10s).

Table 13 Effect of Different Compatibilities on NBR-PP Comuositions

Composition number 4

Recipe PP Chlorine-modified PP ATBN masterbatch 50 MaleicPTETA-modified PP masterbatch NBR Carboxylated copolymer block Nylon-6IBR NBR NBRPP copolymer block resin Phenolic 2.5 3.75 3.75 3.75 dihydrate dichloride Tin(I1) Properties UTS (MPa) Stress at 100% strain (MPa) Young’s modulus (MPa) Ultimate elongation(%) Tension set(%) True stress at break (MPa) Source: Ref. 35.

3

5

1

2

50 0 50

0

50

0

50

50

0 0 0

0 0 0

0 0 0 0

0 0 50 0

0

0 0

0 0 0

0 0

0 0 0 50 0

0

0

50 0

3.75 0.5

0.5

0.5

0.5

8.5

11.9 11.8 239

-

8.2

19.2 14.2 137 250 39 67

23 1 8

9

1 6 0 46 31

180 13

9

0.5

8.7

8

6

7

45 0 0 0

50

50

10 45 0 3.75 2.5 2.5

233 19

10.1 10 170 140

10

24

-

-

0 0 0 0 0 50 0

3.75 0.5 8.8

209 19

-

10

49 0 0 0 0 0 50

1 .o

0.5

18.3 12.7 195 330 43 78

647

Compatibilization of Thermoplastic Blends

glycidyl methacrylate (GMA), 2-hydroxyethyl methacrylate (HEMA), 2-hydroxypropyl methacrylate (HPMA), t-butylaminoethyl methacrylate (TBAEMA), dimethylaminoethyl methacrylate, and 2-isopropenyl-2oxazoline (IPO)were used as the modifiers. It was found that IPO and GMA are effective in compatibilizing the PP-NBR blends [43].The compatibilization of NBR-PP [44]and NBR-HDPE [45] blends has been reported by Thomas andcoworkers. High-impact PScan be obtained from PS and EPDM by the coupling of EPDM and PS in the mixing chamber of a HAAKE plastograph. Lewis acids were addedto the melt, and it was found that rubber became crosslinked or was coupled with the PS molecules and improvement in mechanical properties was observed [46]. Santra et al. [47] reported the in situ compatibilization of low-density polyethylene and polydimethyl siloxane rubber blends using ethylene-methyl acrylate copolymer as a chemical compatibilizer.Ethylene methacrylate (EMA)reacted with the rubber to form EMA-grafted rubber during the melt mixing, whichacts as the compatibilizer. They have conducted dynamic mechanical analysis, adhesion studies, and phase morphology and foundthat 6 wt% ofthe compatibilizer was the optimum quantity required. Els and McGill [48] reported the action of maleic anhydrideon polypropylene-polyisoprene blends. A graft copolymer was found in situ through the modifier, which later enhanced the overall performance of the blend. Scott and Macosko [491 studied the reactive and nonreactive compatibilization of nylon-ethylenepropylene rubber blends. The nonreactive polyamideethylene propyleneblendsshowed poor interfacial adhesion between the phases. The reactive polyamideethylene propylene-maleic anhydride modifiedblends showed excellentadhesion and much smaller dispersed phase domain size. Greco et al. [50] studied the effect of the reactive compatibilization technique in ethylene propylene rubber-polyamide-6 blends. Binary blends of polyamide6-ethylene propylene rubber (EPR) and a ternary blend of polyamide-6-EPR-EPR-g-succinic anhydride were prepared by the melt mixingtechnique, and the influence of the degree of grafting of (EPR-g-SA) on morphology and mechanical properties of the blends was studied.

VIII. CONCLUSION Thermoplastic elastomers are materials that have the properties of vulcanized rubbers but can be processed by techniques associated with thermoplastics. The commercial importanceof TPEs is due to their superior processing properties and economic advantages over conventional rubbers and plastics. TPEs fromrubberplastic blends becameimportant because they combine the superior processability of thermoplastics and the

very good mechanical properties of the elastomers. Both natural and synthetic elastomers can be blended with thermoplastics either by melt blending or by the solution-castingtechnique for the preparation of TPEs. Most of these TPEs are immiscible and incompatible andexhibit poorproperties. This problem can be alleviatedby the proper addition of a suitable compatibilizer, which may be a block copolymeror graft copolymer or homopolymer or low-molecularweightmaterials. The dynamic vulcanization of the rubber phase is widely used as a means to improve the compatibility. There are many structural andtechnical parameters that control the overall performanceof the compatibilized blendsystem. A few theoretical approaches were also discussed. The changes in mechanical and morphologicalproperties of the blends uponthe addition of the copolymer were discussed in detail. REFERENCES 1. M. Xanthos, Polym. Eng. Sci., 28: 1392 (1988). 2. D. R.Pauland J. W. Barlow, J . Macromol. Sci. Rev. Macromol. Chem., C18(1): 109 (1980). 3. D.R. Paul, C. E. Vinson, and C. E. Locke, Polym. Eng. Sci., 12: 157 (1972). 4. A. J. Yu, “Multicomponent Polymer Systems” (N. A.J. Platzer, ed.), Advanced Chemistry Series, Vol. 9 9 , p. 2, Amer. Chem. Soc., Washington, DC (1971). 5. N. G. Gaylord, “Copolymers, Polyblends and Composites” (N.A. J. Platzer, ed.), Advanced Chemistry Series, Vol. 142, p. 76, Amer. Chem. Soc., Washington,DC (1975). 6. C. S. Tucker and R. J. Nicholas, S. P. E.ANTEC Tech. Papers, 33: 117 (1987). 7. M. Hajian, C. Sadrmohaghegh, andG. Scott,Eur. Polym. J., 20: 135 (1984). 8. P. G . Anderson, U.S. Patent, 4, 476, 283 (1984). 9. F. Ide and A. Hasegawa, J . Appl. Polym. Sci., 18: 963 (1974). 10. S. Thomas and R. E. Prud’homme, Polymer, 33: 4260 (1992). 11. Z. Oommenand S . Thomas, Polym. Eng. Sci. 36: 151 (1996). 12. R. Asaletha, M. G. Kumaran,and S. Thomas, Rubber Chem. Technol. 68: 671 (1995). 13. G. Riess, J. Kohler, C.Tournut, andA. Banderet, Makromol. Chem., 101: 58 (1%7). 14. G . Riess and Y. Jolivet, “Copolymers, Polyblends and Composites, (N. A. J. Platzer, ed.), Advanced Chemistry Series, Vol. 142, p. 243, Amer. Chem. Soc., Washington, DC (1975). 15. R. Fayt,R. Jerome,andPh. Teyssie, J . Polym. Sci. Polym. Phys. Edn., 20: 2209 (1982). 16. G. Riess, P. Bahadur, and G. Hurtrez, Encyclopedia of Poly. Sci. Eng. 2, 234 (1985), Second Edn., John Wiley and Sons, New York. 17. L. Leibler, Makromol. Chem. Symp., 16(1): 17 (1988). 18. J. Noolandi and K. M. Hong, Macromolecules, IS: 482 (1982). 19. J. Noolandi, Polym. Eng. Sci., 24: 70 (1984).

648 19a. S. H. Arastasiadis, I. Gancarz,and J. T. Koberstein, Macromolecules 22: 1149 (1989).

20. C. K. Perk, C. S. Ha, J. K. Lee, and W. J. Cho, J . Appl. Polym. Sci., 50: 1239 (1993). 21. Z. Oommen and S. Thomas, Polymer Bulletin, 31: 623 (1993). 22. R. Asaletha, M. G. Kumaran,and S. Thomas, Polym. Plast. Technol. Engg., 34: 633 (1995). 23. R. Fayt R. Jerome, and Ph. Teyssie, Polym. Eng. Sci., 27: 328 (1987). 24. L. H. Chu, S. H. Guo, and H. C. Tseng, J . Appl. Polym. Sci., 49: 179 (1993). 25. T. Inoue, J . Appl. Polym. Sci., 54: 723 (1994). 26. T. Inoue and T. Suzuki, J . Appl. Polym. Sci., 56: 1 1 13 ( 1995). 27. Z. Krulis, I. Fortelny,and J. Kovar, Collect.Czech. Chem. Commun., 58: 2642 (1993). 28. A. J. Moffett and M. E. J. Dekkers, Polym. Eng. Sci., 32: l(1992). 29. C. S. Ha, J . Appl. Polym. Sci., 37: 317 (1989). 30. N. R. Choudhuary, P. P. De, and A. K. Bhowmick, Thermoplastic Elastomers from Rubber-Plastic Blends, Ellis Horwood, England, Chap. 3, p. 79 (1990). 31. Y. Kim, C. S. Ha, T. Kang, Y. Kim,andW. Cho, J . Appl. Polym. Sci., 51: 1453 (1994). 32. D. G. Peiffer and M. Rabeony, J . Appl. Polym. Sci., 51: 1283 (1994). 33. A. Y. Coran and R. Patel, Rubber Chem. Technol., 54: 892 (1981). 34. A. Y. Coran and R. Patel, Rubber Chem. Technol., 55: 116 (1982). 35. A. Y. Coran and R. Patel, Rubber Chem. Technol., 55: 1063 (1982).

Asaletha et al. 36. G . Riess, P. Bahadur, and G. Hurtrez, Encyclopedia of Poly. Sci. Eng. 2, 234 (1985). Second Edn., John Wiley and Sons, New York. 37. M. Frounchi and R. P. Burford, Iranian J . Polym. Sci. Technol., 2 : 59 (1993). 38. D. J. Elliot, Thermoplastic Elastomers from Rubber-PlasticBlends, Ellis Horwood,England,Chap. 4, p. 121 (1990). 39. Chuan Qin,J.Yin,and B. Huang, Polymer, 31: 663 (1990). 40. C. P. Rader and S. A. Sabet, Thermoplastic Elastomers from Rubber-PlasticBlends, Ellis Horwood,England, Chap. 6, p. 177 (1990). 41. Y. Wang and S. Chen, Polym.Eng.Sci., 21: 47 (1981). 42. K. E. George, Rani Joseph, D. Joseph Francis, and K. T. Thomas, Polym. Eng. Sci., 27: 1137 (1987). 43. M. C. Liu, H. Q. Xie, and W. E. Baker, Polymer, 34: 4680 (1 993). 44. S. George, R. Joseph, K. T. Varughese, and S. Thomas, Polymer 36: 4405 (1996). 45. J. George, R. Joseph, K. T. Varughese, and S. Thomas, J . Appl. Polym. Sci.,57: 449 (1995). 46. E. Mori, B. Pukanszky, T. Kelen, andF. Tudos, Polymer Bulletin, 12: 157 (1984). 47. R. N. Santra, B. K. Samantaray, A. K. Bhowmick, and G . B. Nando, J . Appl. Polym. Sci., 49: 1145 (1993). 48. C. EISand W. J. McGill, Plast. Rubb. Comp. Proc. Appl., 21: 115 (1994). 49. C. E. Scott and C.W. Macosko, International Polym. Proc., 10: 1 (1995). 50. R. Greco, M. Malinconico, E. M. Celli, G . Ragosta, and G . Scarinzi, Polymer, 28: 1185 (1987).

43 Polymer Blends and Alloys B. S. Kaith, A. S. Singha, and Sunil Regional Engineering College, Hamirpur, India

1.

a multicomponent system form the basis of the polymer blending [26-281. Most materials, be they natural or synthetic, have limBlending of polymers is the most versatile way for ited utility. However, technical ingenuity has increased producing new materialswithtailor-made properties, the utility of these materials beyond anyone’s wildest which are not possible via copolymerization, for deimagination. The enormous range of steel that can be manding applicationsin the polymer technology. Block produced by adding carbon or other elements to give it and graft copolymerization share manycommon feathe required balanceof properties, such as strength and tures and purposes as blends, but these materials differ hardness, related to changes in their microstructure from blendsby only a few chemicalbonds. In block and [l-31 is just one example. graft copolymers, the thermodynamics of phase separaAn analogy can be drawn in the case of pure polytion of the constituent polymers is different from that mers, i.e., on their own they exhibit extremely limited of common polymer blends. The existence of covalent utility. In fact, many pure polymers are difficult to probonds greatly influence or control the phase behavior cess into satisfactory products. In order to make them and improvethe miscibility of two dissimilar polymeric useful, certain additives are incorporated to give ease of phases, thereby altering and optimizing their physical processibility, and useful mechanical and load-bearing and mechanical properties. However, in PBAs proper properties. Based on the broad principles governingthe selection of blend ratios, processing techniques, and use science and technology of compounding, the skill and of a selected compatibilizer may be tried to derive differexperience of the compounder plays a vital rolein delivent levels of improved physical and mechanicalproperering the rightmaterial for a particular application ties; selective chemical modifications may also offer en[4-251. hanced property balance making in tailor-made One of the major areas of thrust in science and techpolymers of special significance [29-381. nology for the decade has been the field of materials, of In a fundamental sense, the miscibility, adhesion, which polymer blends and alloys (PBAs) havethe major interfacial energies, and morphology developed are all share and are one of the most dynamic sectors in the thermodynamically interrelated in a complex wayto the polymer industry. Considering the worldwide scenario interaction forces between the polymers. Miscibility of in the field of PBAs andtheir growing commercial appli- a polymer blend containing two polymers depends on cability, this field is probably the right and timelychoice the mutual solubilityof the polymeric components. The for raw material and polymer educators, scientists, and blend is termed “compatible” when the solubility patechnologists. PBAsare gaining importance in academic rameter of the two components are close to each other and industrialresearch. These blends have some unique and show a single-phase transition temperature. Howproperties that are different from the basic polymers ever, most polymer pairs tend to be immiscible due to from which these have been produced. To improve the differences in their viscoelastic properties, surface-tenprocessing behavior for end use, one polymer blending sions, and intermolecular interactions. According to the with another polymer isa common practice. The exploiterminology, the polymerpairs are incompatibleand tation of certain unique sets of properties of the individshow separate glass transitions. For many purposes, ual polymer for the benefit of the overall properties of miscibility in polymer blends is neither required nor deINTRODUCTION

649

650

Kaith et al.

sirable. However, adhesion between the components is one of the most important factors. Optimum adhesion between the dispersed phase (domain) andthe matrix are gainfully attained through the use of a third component known as the reactive compatibilizer,which has a chemical structure matching those of the component polymers [9,10,13,39-531. Following are some of the important terms used in the field of polymer blends and alloys [54].

cally distinct and mechanicallyseparable materials. by mixingthe separate 2. Composites can be made materials in such a way that the dispersion of one material in the other can be done in a controlledway to achieve optimum properties [10-221. 3. The properties are superior and possibly unique in some specific respects to the properties of the individual components.

1. Polymer Blends: Mixture of chemically differ-

2. 3.

4. 5.

6.

7.

8.

ent polymers or copolymers with no covalent bonding between them. Polymer Alloys:A class of polymer blends, heterogeneous in nature with modified, controlled interfacial properties or morphology. HomologousPolymerBlends: A subclass of polymer blends limited to mixtures of chemically identical polymers differing in molecular mass. Miscible Polymer Blends: A subclass of polymer bends encompassing those blends that exhibit single-phase behavior. Immiscible Polymer Blends: A subclass of polymer blends referring to those blends that exhibit two or more phases at all compositionsand temperatures. Partially Miscible Polymer Blends: A subclass of polymer blends including those blends that exhibit a “window” of miscibility, i.e., they are miscible onlyat certain concentrations and temperatures. Compatible Polymer Blends:A term indicating commercially usefulmaterials, mixture of polymers with strong repulsive forces that is homogeneous to the eye. Interpenetrating Polymer Network (IPN): A subclass of PBs reserved for the mixture of two polymers where both components form a continuous phase and at least one is synthesized or crosslinked in the presence of the other [8].

The PBAs [S], however,must be distinguished from polymericcomponents that aredefined as follows: 1. Compositesconsisting of two or morephysi-

II. INTERNATIONALSCENARIO

The world production of plastics in 1995 is projected at 76 million metric tons (mT) with an annual growth rate (AGR) of3.7%. The expected AGR of PBAs is 12% and that of composites 16%. In 1987,21% of polymers were used in blends and 29% in composites and filledplastics [56]. If this trend continues, by 1995 all manufactured resins will be used in multiphase polymeric systems, Two factors moderating the tendency are: 1. A need for single-phase polymersin some applications, e.g., polytetrafluoroethylene, VHMWPE, light-sensitive polymers, etc. 2. Use of polymer alloys and blends as matrices for composites, reinforcements, and foams.

There has been a rapid growth of the demand for plastics fromless than 20 billion pounds in 1970 to nearly 50 billion pounds consumed in the United States in 1986, mostly due to the substitution of traditional raw materials. All over the world, plastics have replaced metals, glass, ceramics, wood papers, and natural fibers in a widevariety of industries includingpackaging,consumer products, automobiles, buildingand construction, electronics and electrical equipment, appliances, furniture, piping, and heavy industrial equipment [57-1211. Consumption patterns of PBAs in some countries are shown in Tables 1and 2. The market study conducted by Frost and Sullivan [l221 forecasted that sales of PBAswould continue strongly in the United States with consumption growing from the 1990s total of 363,000 MT to 466,000 MT in 1995.

Table 1 PBA Global Consumption Patterns (in kMT)

USA 1996

PBAs 1986

1996

High-volume PBAs Engineering plastics (PBAs) Speciality plastics (PBAs) Growth rate

Japan

W. Europe

1986 76109 179 105 8

270 216 409 36 9.5%

23

141 225 1

88 0.5

8.2%

7

-

10.3%

Polymer Blends and Alloys

651

Table 2 Consumption of Six Main Families of PBAs

added value for the recycling of plastics, thus, itis an environment-friendlymechanism for plastic wastes management.

in Western Europe (MT) PP-EPDM: PC-ABS: PS-PPE: PC-PBT: Nylons: PVC-ABS:

111.

90,000 49,000 43,000 31,000 18,000 16,000

Automobile applications: Appliances: Electronics: uses: Industrial Construction: Consumer applications:

56% 19% 15% 10% 7% 4%

ADVANTAGES OF BLENDING

The combination of two or more commercially available polymers through alloyingor blending represents an inexpensive route to product differentiation for suppliers. For a processor and end user, alloying and blending technology permits tailoringof a polymer compoundto their specific applicationrequirements, often at a lower cost than the current material and over a shorter developmental period. Alloy and blend development is typically market driven and requires an ongoing dialogue between supplier and customer to enable commercial success [123]. The major advantages of blending are as follows: 1. High-impact strength 2. Easy in processibility 3. High-tensile strength 4. Modulus/Rigidity 5. Head deflection temperature 6. Flammability 7. Solvent resistance 8. Thermal stability 9. Dimensional stability 10. Elongation 11. Gloss

There are number of reasons that justify the positive prospects for not onlythe established PBAs butalso for the new ones: 1. PBAs close technological and economic gaps

2.

3. 4.

5. 6.

IV. SELECTION OF BLENDING COMPONENTS Selection of blending components is a difficult but important task for polymer technologists.This task can be achieved by the selection of blending components in such a way that the principal advantages of the first polymer willcompensate for deficiencies of the second polymer and vice-versa. At the beginning, it is essential to highlight the advantages and disadvantages of the polymers (Table3)that are tobe used andthen make combinations to suit the requirements for a particular set of properties. All this should be done while keeping in mind the processibility of the materials, the limitation of the equipment, andthe stability of the materials. Moreover, in the selection of the blend components, the economic factor must also be taken into consideration. This can be illustrated by the following examples. A.GolfBalls Originallygolfballs were madebystuffing feathers under high pressure into a stitched leather container. Subsequently,gutta-percha or balata was used. With the passage of time, the use of gutta-percha became obsolete. Around 1900, a ball witha core made from winding rubber threads was prepared. However, it was a tedious process and the winding techniques differed. With advancementsin technology, this process has beenre-

Table 3 Formulation of a “0”Type of Mount

in a Load Deflection Operation

Phr

NR

Bromobutyle rubber among the thermoplastics. ZnO Acid The production of PBAs meansthat a successful Polybutadiene nitrile-acrylate combination of properties can be achieved to IIPD almost perfectly fitthe respective requirements FEF of an application. Black pine tar The development time neededfor PBAs before Chlorobutyl-costyrene a product can be launched into the market is Sulphur shorter when compared with new materials. TMTD The blends can be manufactured without high Physical Properties additional capital investment. There is familiarHardness ity with the processing of such materials. Tensile strength (MPa) Comprehensive know-how in PBAs technology Mod. 300% MPa has been developed andconstantly is expanded. Percent elongation PBAs technology provides an opportunity for

1nn.n

75 .0 ” ”

0.0

75.0

5.0 1.S 1.o 1.0 10.0 10.0 1.5 1.S

0.2

5.0 1.S 1.o 1.o 10.0 10.0 1.5 1.5 0.2

40.0-41.0 14.4 4.9 650.0

40.0 12.7 5.0 630.0

652

Kaith et al.

Table 4 TypicalThreadRecipe

Percentage weight 100.0

cis-Polybutadiene Aldehydelamine Zinc 4,4‘ Dithiodimorpholine Sulphur

0.5

3.0 1.O 1.o

Cure time at 135°C is 2 h.

placed bya different methodof preparing golf balls, i.e., a two-piece ball. One the is “inner core” and the other is the “outer cover”; both are made ofseparate materials.

l . Inner Core Since 1960, the inner core has been made from cis-polybutadiene by the compression moulding technique. This replaced the earlier material made froma suspension of barytes or bentonite clay in water and glycerine or the winding of rubber threads made from cis-polyisoprene, either from latex or a dry rubber compound. A typical thread recipe is given Table 4. This recipe gives an elongation of about 100% with high-tensile strength. But this threaded core has been replaced by a solid mouldedcore, which is based onthe use of a crosslinkage system. The monomers preferred are metal salts of polymerizable organic acids, such as zinc diacrylate and zinc dimethacrylate. These may be formed separately or in situ during compounding. Generally, metal-containingpolymerizablemonomers produce cores of higher resilience, for which an extra protection cover is needed. A typical formula is shown in Table 5.

cover material ( 16OoC), separate moulding procedures are adopted for the cover and the inner core. Technology Cover compounds with desired properties are obtained by blending twoor more differentionomers or by blending withother thermoplastic materials. Half shells produced by injection mouldingare linked together for ease of handling. Since ionomersrequire high-moulding temperatures ( 160”C), corelcover assemblies can be put into a hot mould, closed rapidly, and then immediately cooled. Special mould coatings are necessary for good release property. Finally, the equator sprue line is, accurately trimmed after moulding. In a particular application involving a typical “0” type mount, the nitrile rubber (NR) compoundcauses a resonance frequency of 28 Hz, whereas resonance was previously tolerable only well below 20 Hz. In thiscase, a blend of NR and bromobutyl rubber is more suitable. A comparison is shown in Table 3. Similarly, in the case of a cup type mount, an initial attempt to restrict resonance amplification featured the use of white factice in a NR formulation. However, the successful formulation was a blend of predominantly bromobutyl with NR, as shown in Table 6. B. Automobile Panels and Bumpers Generally, 85-90% of rubber modified polypropylenes (PP) are used in the automobile industry. It allows for a better way of reducing the weight of the vehicle, increased longevity, improved impactresistance, a glossy appearance, and reduction in cost. Suitable blends contain 10-35%, by weight, of ethylene propylene diene rubber (EPDM). Their specificgravityvariesfrom 0.87-0.89 g/cm3. Impact strength is retained at normal temperatures down to - 30°C and in stable to ultraviolet radiationsand heat aging. Recently, the blends have

2. Cover Material Since 1960, gutta-percha or synthetic balata has been replaced by “ionomer” for use as a cover material in golf balls. It is based on a combination of ethylene and a carboxylic acid. It can be processed at a reasonable temperature to produce good, tough ball covers. Because of higher processing temperature of the ionomer

Table 5 Formula for Higher Resilience Cores

Percentage weight cis-Polybutadiene Zinc diacrylate Zinc Dicumyl peroxide (40% active) Cure time at 160°C is 20 min.

100.0 35.0 18.0

5.0

Table 6 Formulas for a Cup Type Mount

Phr NR

Bromobutyl rubber FEF Black pine tar White factice Chlorobutyl-costyrene Sulphur Physical Properties: Hardness Tensile strength (MPa) Mod. 300% MPa Percent elongation Amplification at resonance

100.0 0.0

1.5 1.5

25.0 75.O 5.0 10.0 20.0 1.5 1.5

35.0 13.3 4.5 650.0 8.5

35.0 11.1 4.7 675.0 3.5

5.0 10.0

20.0


been reinforced withfibrous fillers, such as carbon and glass fibers, for higher impactstrength. The PP-EPDMS blends have been replaced by other polymer blends such as poly(2,6 dimethy1:4 phenylene ether)-polybutylene terephthalate (PPEPBT), polycarbonates-styrenecomaleic-anhydride (PUSMA), poly(2,6 dimethyl 1:4 phenyleneether)-polyamide(PPEIPA),andacrylonitrile butadiene styrene-polycarbonates (ABSPC), etc. [W. V. CLASSIFIED PBAs

A.ThermoplasticElastomers(TPEs) TPEs are a new class of materials, combining the processibility ease of thermoplastics and the functional performance of conventional thermoset elastomers. TPEs need no vulcanization and canbe processed using conventional techniques, such as injection moulding, blow moulding, extrusion, etc. The morphology of the TPE provides unique properties to it, and it is a phase-separated system. The first phase is hard and solid, while the second one is soft andrubbery at room temperature. The properties depend on the nature and amount of the hard phase present. The temperature range of applications is determined bythe combined effect of the glasstransition temperature of the soft phase and the melting point of the hard phase. There are two categories of TPEs: (1) Block copolymers of two or more monomers, and (2) Simple blendsof thermoplastics and elastomers. Plastics, such as PE, PP, polystyrene (PS), polyester, and nylon,etc., and elastomers such as natural rubber, EPDM, butyl rubber, NR, and styrene butadiene rubber (SBR), etc., are usually used as blend components in making thermoplastic elastomers. Such blends have certain advantages over the other type of TPEs. The desired properties are achieved by suitable elastomers/plastic selection andtheir proportion in the blend. B.ElastomericAlloys

(EAs)

EAs are generated from the synergistic interaction of two or more polymers possessingproperties better than are a special class of thermothose of a simple blend and plastic elastomers. In EAs the elastomer is crosslinked and dispersed in a continuous matrix of thermoplastic under a dynamic condition resultingin a fine dispersion of fully crosslinked rubber particles in EA. Two-phase EAs differ from the common thermoplastic elastomers containing two-phase morphology by: (1) the presence of a high degree of crosslinked elastomer chains, and (2) the degree of interaction between the thermoplastic and elastomeric phases. The alloying process between the two polymers ultimately attains the following characteristics: 1. Better tensile strength. 2. Lesser tension and compression set.

653

3. Greater resistance toward solvents and greater

flex resistance. 4. Retention of properties at elevated temperatures. 5. More consistent processibility. Conventional crosslinking agents, such as sulphur, accelerators, and peroxide, etc., used in dynamic vulcanization and melt-mixingof two polymers is the first step in the preparation of an EA. The dynamic vulcanization is done by dynamic shear at a high temperature to activate the process. A compatibilizer is sometimes used to overcome the interfacial tensionbetween the two phases of dissimilar polymers. It enables a fine dispersion of highly crosslinked rubber particles. The function of the compatibilizer isto provide greater, but not total, thermodynamic compatibility between the two polymers [8]. VI. PREPARATION OF PBAs

PBAs are generally prepared by three commercial methods [57,124]: latex mixing,solutionmixing,andmelt mixing. A.Latex

Mixing

Latex mixing is the most important technique for the preparation of commercial PBAs. In this process two important factors are to be taken into consideration:

1. pH of latex: It should be approximately equal for better mixing of two polymers. 2. particle size: If small particle size latex is to be blended with a latex of large particle size, stabilizer shouldbe added to the latex with the larger particlesize prior to blending the two latexes. Polymers are suspended as microparticles in the latex and interactions between these microparticles are prevented by the presence of adsorbed suspendingagent and soap molecules. Blending results in a random suspension of dissimilar particles in the mixture of latexes, each unaffected by the other. Rate of flocculation depends entirely on the stabilizer and not on the polymer characteristics as such. Coagulated masscontains an intimate mixtureof the polymers. Acrylonitrilebutadiene styrene (ABS) polymers[23-251 may be prepared by this method. Compounding and palletizing of the latex blended materialissometimes done bymeltmixing.In such cases precautions must be taken to avoid thermal and shear degradation of the blended polymer.

B. SolutionMixing This method is used for polymers when they are not amenable to melt processing.In this technique, a diluent

Kaith et al.

654

is added in order to lower the temperature and shearforce requirement for satisfactory mixing of polymers without anydegradation. However, attempts to remove the diluent fromthe mixture may leadto a change in the domain sizes in the blend, which in severe cases may cause gross-phase separation [ 124,1251. One of the most important solution blend polymers is high-styrene resin, which is manufacturedby several companies worldwide.This is a latex blend of high-styrene rubber and normal styrene butadiene rubber. The different high-styrenemaster batches are available in the world as:

roduct nameCompany ~

1. Japan synthetic rubber 2. Kubochem 3. BF Goodrich Tire Group Ameripol Tire Div. 4. Chemische Werke Huels A.G. 5. Nippon Zeon Co. Ltd., Japan 6. Synthetics & Chemicals Ltd., India 7 . Apar Ltd., India 8. Apar Ltd., India

SBR 0051,0061 AF POL 537,539 Ameripol1901,1904 DURANIT B NIPOL HS 850, HS 860 SYNAPRENE 1958 POWERENE 958 APARENE198

C. MeltMixing In this technique, the melting ofconstituent polymers is of utmost importance for the generation of a uniformphase morphology, which finally controls the performance of the blended materials.In general, melt mixing is used for a system in which thermal degradation does not ordinarily occur. To prepare the commercial PBAs, uniform blending of immiscible polymer pairs is often done through the use of a reactive compatibilizer. This method involvesthe preparation of a master batch of the compatibilizer and the dispersed phase, which is then blended with the matrix through the melt-mixing technique. In this mixing process, contaminants such as solvent and/or diluents as well as their removal problems can be avoided. Degradation of the polymers is avoided by proper maintenanceof the viscosity and shearing rates. Other methodsof blending include: (1)fine powder mixing, and (2) monomer as a solvent for other components of the blend, followed by polymerization for making an interpenetrating network (IPN) [15]. VII. COMPATlBlLlZATlON OF PBAs There are many definitionsof polymer compatibility. On one hand, compatible polymers are the polymer mixtures that have desirable physical properties when

blended, while, on the other hand, those mixtures that do not exhibit gross symptoms of phase separation when blended are said to be compatible blends.However, the most used definition of polymer compatibility the total is miscibility at the molecular level of blending components. Only a few blends form complete miscible systems, which are characterized by a single T, and homogeneity of 5- to 10-nm scale. Most of the blends are immiscible, possessing a phase-separated morphology. In such cases, important physical and mechanical properties are related to the finely dispersed phase and resistance to gross-phase separation. Moreover, it isreasonable to believe that PBAs exhibiting no gross symptoms of phase separation on blending and having desirable properties show at least some mixing of polymer segments ona microscopic scale. This implies the establishment of some degreeof thermodynamic compatibilityor the existence of physical constraints such as crosslinking, graft or block copolymer sequences, which are expected of a single-phase material. The following exampleprovides a clear understanding of the term miscibility. If one refluxes PS on a mill with polyphenyl oxide (PPO) as thesecond component, a single-phase polymersystem results that is thermodynamically stable and is an example of a miscible blend, i.e.,polyvinyl chloride (PVC)-NBR system. (already commercialized under the tradename of Noryl by General Electric Co., USA. It has been observed that if PS is refluxed with an equal amount of PVME on a mill at 80"C, a clear one-phase mixture results, indicating a miscible system at that temperature. However, if the temperature of the system is raised to 14O-15O0C, two phases appear in the system, indicating the development of immiscibility. Thus, the driving forces for the transition fromthe miscible state (single phase) to immiscible state (two phase) are thermodynamic in origin and do not depend on the extent or intensity of mixing. As a result of recent advances in the field ofcontinuous reactive processing, particularlyin extrusion compatibilization of polymer blends through reactions during compoundingisbecomingincreasingly important. The versatility of twin-screw extruders as the reactor combined with wider availabilityof polymers functionalized with a variety of reactive groups appear to be among the reasons contributing to the growth in the field of reactive compatibilization [126,1271. Methods depending on reactive compatibilization can be classified[ 1261 as follows: 1. Those involving the use of functionalized blend componentsthat produce in situ compatibilizing interchain copolymers: some commonreactive functionalities includecarboxyl, epoxy, isocyanate, and anhydride, etc. 2. Methods involving the addition of a third highmolecular weight component capable of reacting with at least one blend component. In this


case, compatibilization may result through covalent or ionic bonding. 3. Through the addition of a low-molecular weight component that may promote copolymerformation and crosslinking or co-crosslinking reactions. Blend of (1) and (2) type categories mostly include the modification of engineeringthermoplastics with another thermoplastic or rubber. PS-EPDM blends using a low-molecular weight compound(catalyst) Lewis acid have been developed [126]. Plastic-plastic blends, alloys of industrial importance, thermoplastic elastomers madebydynamicvulcanization,and rubber-rubber blends are produced by this method.

VIII. CHARACTERIZATION OF PBAs Complexitywithinhomopolymers as well as that of PBAs have made the task of analysis and characterization a difficult one. Basically, the task of analysis and characterization of PBAs is not different from that of simple low-molecular weight polymers, provided adequate solubility andsites are available for accepting artificial stimulationresponses to those stimuli that may be used as functional tools for characterization. Properties of the blendmainlydependon the homogeneity of blends. The processes that are used for characterization of the PBAs are discussed in the following sections [128-1311. Microscopy A. Electron and optical microscopesare being used to see blend homogeneity. Elastomer-plastic blendsare somewhat easier to identify than elastomer-elastomer blends because normal staining techniques, e.g., osmium tetraoxide, can be used in the case of plastic-elastomer blends. Normally, there are two methods that are followed for examining the blend surface by electron microscopy.

I . EboniteMethod The blends are cured with a sulphur-sulphonamide zinc stearate mixture so that the elastomer phase is hardened and then microtomed. 2. CryogenicMethod The blends are frozen in liquid nitrogen and then microtomed and stained with osmium tetraoxide, which stains only unsaturated elastomers.

B.SolubilityMethod Solubility differences among components of polymer blends can be utilized in different ways to identify the polymer blends.

655

C.OpticalProperties Light scattering of polymer blendsalso helps incharacterizing the different phases of PBAs.

D. ThermalandThermomechanical Analysis Thermal analysis helps in measuring the various physical properties of the polymers. Inthis technique, a polymer sample is subjected to a controlled temperature program in a specific atmosphere and properties are measured as a function of temperature. The controlled temperature program may involve either isothermal or linear rise or fall of temperature. The most commonthermoanalytical techniques are: (1) differential scanninganalysis (DSc), (2) thermomechanical analysis (TMA), and (3)thermogravimetry (TG). These techniques help in providingthe following information: specificheat, enthalpy changes, heat of transformation, crystallinity, melting behavior, evaporation, sublimation,glass transition, thermal decomposition, depolymerization,thermal stability, content analysis, chemical reactions/polymerization linear expansion, coefficient, and Young's modulus, etc. DSC helps in determining the glass-transition temperature, vulcanization,and oxidative stability. TG mainly is applied for the quantitative determination of majorcomponents of a polymersample.TMA or DLTMA(dynamicloadthermomechanicalanalysis) measures the elastic properties viz. modulus. Thermal analysisof homopolymer samplesare simpler than those of blends. Separate thermal analysis of individual polymer components are made before doing the same for a blend in order to get more accurate and proper information on thermalcharacteristics. It is possible to distinguish between SBR andbutyl rubber (BR), NR and isoprene rubber (IR) in a vulcanizate by enthalpy determination. In plastic-elastomer blends, the existence of high Tgand lowT, components eases the problems of experimental differentiation by different types of thermal methods. For a compatible blend, even though the component polymershave different T, values, sometimes a single TBis observed, which may be verified withthe help of the following equation: " "

where w1 and TB!, w2 and Tm are the weight fractions and the glass-transition temperature of the two polymer components 1 and 2,respectively [129]. Glass-transition temperature of the blend can also be determined by the differential scanningcalorimeter, which gives indication about the nature of the blend.

v,)

Kaith et al.

656

E.

IR and NMR Spectroscopy

In multiphase polymericsystems, the properties of the end products do not solely depend on the properties of the pure components, but other various parameters also havea great impact (Fig.‘l). In order to emphasize these factors, the following systems are taken into consideration: (1) elastomer toughened styrene system, (2) elastomer toughened polycarbonate blends, and (3) direct reactive blend processing.

The miscibility of two or more polymers in the solid state can be investigated by this technique.

F. X-Ray Diffraction Method For the structure evaluation this method is one of the most suited methods. G.

A.

GelPermeationChromatography

In the late 1940s,the demand for styrene homopolymers (PS) and styrene-acrylonitrile copolymers (SAN) was drasticallyreduced due to their inherent brittleness. Thus, the interest was shiftedto multiphase high-impact polystyrene (HIPS) and rubber-modified SAN (ABS). In principle, both HIPS and ABS can be manufactured by either “bulk” or “emulsion” techniques. However, in actual practice, HIPS is made only by the bulk process, whereas ABS is produced by both methods [132,133].

It is basically a fractionation process that depends not only on molecular size, but also on chemical composition, stereo-configuration, branching, and crosslinking. For multicomponent systems, fractionation with different ion polymolecularity, chemical heterogeneityand sequence length distribution, solubility or elution fractionation is of primary importance. Therefore, gel permeation chromatography or size exclusion chromatography is usedas an important tool for the characterization of PBAs. IX.

Elastomer Toughened Styrene System

l . Bulk Process Initially, this process consists of a rubber dissolved in a monomer(s). During polymerization,the phase homogeneous system transforms into the heterogeneous system over a wide conversion range. The heterogeneous phase originates due to the formation of a subemulsion [ 1341consisting of polystyrene droplets in a rubber solution and can be correlated with the viscosity conversion profile [135].

BLENDMORPHOLOGY/PROPERTY RELATIONSHIP

One of the basic criteria for the process choice is the optimum size and shape of the dispersed phase leading to better overall properties. As a brief guideline, Table 7 canbeused to choose the correct manufacturing process.

Table 7 Effect of Morphological and Molecular Aspects on the Technological Properties

aspects Molecular Morphological aspects

sphase) (dispersed properties Technological Variation due to increasing

sharply

Reactive Molecular Processing Phase-volume weight and size Particle distribution distribution compatibilization its aids its ratioand Increases Impact Increases, Increases, reaches Increases reaches Increases sharply strength Modulus

Flow nally inferior properties es ains Initially s Enhances Enhances crazing Solvent resistance

Decreases sharply Becomes

followed by a decrease Decreases slowly

followed by a decrease Increases slowly

Improves

Improves

deteriorates Remains unchanged remains deflection Remains Initially Lowers Lowers Heat unchanged temperature by an increase

Does not change slowly Becomes inferior

Decreases Improves


657

P R O D U C TP R O P E R T I E S (STIFFNESS,TOUGHINESS, FLOWABILITY, HEAT DISTORTION TEMPERATURE, STRESS CRACK RESISTANCE, GLOSS, TRANSPARENCY, WEATHER RESISTANCE,CREEP etc.)

MORF'HOLOGICAL PARAMETERS * PHASE-VOLUME RATIO * PARTICLE SIZE Kc ITS DISTRIBUTION * PARTICLESTRUCTURE * DEGREE OF COMPATIBILITY * REACTIVE COMPATIBILITY e.g. Grafting, Cross-linking etc. AND RELATED ASPECTS

ADDITIVES

COMPONENT (Dispersed) PARAMETERS AND(Continuous)

-1

PROCESSING METHODOLOGY COMPONENT-2

-

MANUFACTURING BULK/EMULSIONetc. METHODOLOGY - MELT/SOLUTIONetc. Figure 1 Molecular and morphological parameters influencing technological properties of multiphasic polymer blends.

Molecular Parameters The molecular weightof the continuous phase is an important parameter that affects the mechanics and the melt flow ofthe end product. It can be controlled by the use of a suitable chain transfer agent (e.g., fer?-dodecyl mercaptan CT = 4.0) or their combinations (e.g., primary mercaptans CT = 26.0 and dimeric a-methyl styrene CT = 0.1) [132]. The particle size of the dispersed phase depends upon the viscosity of the elastomer-monomer solution. Preferably the molecular weight of the polybutadiene elastomer should be around 2 X lo5 and should have reasonable branching to reduce cold flow.Furthermore, the microstructure of the elastomer provides an important contribution toward the low-temperatureimpact behavior of the final product. It should also be emphasized that the use of EPDM rubber [l361 or acrylate rubber [l371 may provide improvedweatherability. It has been observed that with an increase in agitator speed the mean diameter of the dispersed phase (D) decreases, which subsequently levels out at high shear [138-1411. However, reagglomeration may occur in the case of bulk

ABS at high shears [142]. B also increases with the increase of the viscosity ratio of the phases [143]. Any additive, such as different surfactants and grafted copolymers that lower the interfacial tension, decrease the mean particle size of the dispersed phase (Fig. 2). In order to enhance compatibility of the two phases, the length of the block copolymers, acting as a surfactant, must be greater than the critical entangled chain length.In general, particle size < l p m is not effective in toughening.Particle size between 1-2 p m is satisfactory in the toughening process. However, submicron particle size can act as a good toughnerif the dispersed phase is crosslinked. It is interesting to note that blending of two different grades of high-impact polystyrene (HIPS) having widely differentdispersed phase particle sizescaneasilybeused for optimization of impact strength and gloss (Table 8) [143]. Some common structural features of the dispersed phases that can be observed in HIPS are shown in Fig. 3 [144]. The importance of the morphological aspects of the HIPS on their Izod impact strength and gloss has been

0

658

Kaith et al.

PB: 8% BCP 3 0%

PB:7.6%

PB:

BCP : 0.57%

BCP : 1.14%

PB:5.6%

7.2%

BCP

: 3.43%

Figure 2 Influence of added styrene-butadiene block copolymer (BCP) on the particle size.

highlighted for some commercially available HIPS in Fig. 4 [132]. Another problem faced during cooling of the melt is the mismatch of the shrinkage behavior of the continuous and the dispersed phases. Triaxial stress thus developed aroundthe dispersed phase will lead to crazenucleation under load. Void formation withinthe system can be avoided by the incorporation of a minimum amount of crosslinks, which will help in retainingthe elasticity, load modulus, and glass-transition temperature. Those thermally activated crosslinking will significantly affect the swelling index of HIPS-ABS [145]. Morphological Parameters Phase-value ratio depends upon the following factors: dispersed phase contents, number and size occlusion, and degree of reactivity. During the manufacturingprocess, if the grafting increases during early stages of the reaction, the phase volume will also increase, but the size of the particles will remain constant [146-1481. Furthermore, reactor choice plays a decisive role. If the continuous stirred tank reactor (CSTR) is used, little grafting takes place and the occlusion ispoor and, consequently, the robber efficiency is poor. However, in processes akin to the discontinuous system (e.g., towerkascade reactors), the dispersed phase contains a large number of big inclusions. Dispersed Phase Particle Size The size of the dispersed phase is effectively dea small conversion zone (between cided within

phase-inversion and high viscosity). However, within this zone it depends upon the agitator speed, resulting shear stress, viscosity ratio of phases, and nature of the interphase. It can be qualitatively summarized the in fol-

Capsules

Shells

Rods

Rod Clkrters

Droplet8

Droplet Cluctmr

Figure 3 Particle structure in HIPS.

Table 8 Impact Strength and Gloss of HIPS Blends

Rubber particle Large particles Rubber phase Notched impact Izod size (pm) blends in 0.6 96.3 81.3

3.0 0.6

3.0

1.85 1.85 1.90

(%) 0.0 5.0 30.0

volume fraction strength 0.200 0.450 0.213 0.275

ft. Gloss lbs/in. 0.90

100.0 50.0


659

2.8 Particle diameter (um) Izod140 Impact Strength (J/m) Gloss (arbitrary unit)

1.o

5.8 125

I80

18

35

60

(ABS=80)

Particle diameter (um) 76 Strength (J/m) Izod Impact Gloss (arbitrary unit) 80

0.75 140

_-

70

76

0.30 90

(ABS=80)

Figure 4 Correlation between morphology, impact strength, and gloss

lowing equation [149]:

where

D, = critical drop break-up diameter

K, K ' and p = constants 7 7 , and

A = time constant = viscosity of the continuous and dispersed phases, respectively. 7 = shear stress.

vd

Transparent Heterogeneous Blends In general, the multiphasic heterogenous nature of the impact grade styrene-based polymersis the root cause of their opaque-turbid nature. In determiningthe transparency of the blends, size and the size-distribution pattern of the dispersed phase along with the refractive index difference between the continuous and the dispersed phases are two very important criterion [133]. High-Heat Distortion Grade Blends High-distortion temperature (HDT) of the styrenic materials is around 100°C. Blending with polymers of high HDT may lead to blends of high HDT. Both compatible and noncompatible polymers can be used judiciously for this purpose, e.g., PS is compatible with PPE (HDT = 220°C) [l501 on a molecular level. Therefore, it is not surprising that a blend of HIPS with PPE will

of HIPS.

exhibit the same morphology as those of the original HIPS. However, the size of the dispersed phase is smaller thanthe conventionalsystems. Thus, the proper formation willlead to the blends exhibiting high-heat resistance, high-impact strength, high-dimensional stability, high-hydrolytic stability, and halogen-free flame retardant compositions. B.

Elastomer ToughenedPolycarbonate Blends

Heterogeneous compatible blendsof preformed elastomers and brittle plastics are also an important route for the development of blends of enhanced performance with respect to crack or impact resistance. Polycarbonate blends with preformed rubber particles of different sizes have been used to provide an insight into the impact properties and the fracture modes of these toughened materials. Izod impactstrength of the blends having 5-7.5 wt% of rubber particles exhibits best overall product performanceover a wide rangetemperature (RT to -40°C) [151-1541.

C. DirectReactiveBlend Processing Recently, it has been proposedthat direct extrusion processing may be a far better alternative for the blend preparation. It has various advantages, such as a onestep fabrication process, a reduction of the production

660

cost, and it minimizes the thermal degradation of the polymer blend components by reducing their thermal history from compoundingto moulding. Sakai and coworkers [ 1551 have developed a direct compounding and extrusion system, i.e., a hybridized processing system for manufacturing PBAs with excellent impact strength. The main purpose of this direct reactiveblendprocessingtechnologyis reduction of moulding pressure and temperature, utilization of lowmeltviscositythermosetting resins, improvement in physical properties of resultant thermosets, such as impact strength, to obtain blends/alloys equivalent and/or superior to corresponding thermoplastics by incorporation of rubbery. The process is composedof the following three distinct stages:

First Stage Mixing and kneading of the thermosettingresin, rubbery components and other additives.

Kaith et al.

A. Commercial PBAs

Commercial thermoplastics are the engineering materithat als containing twoor more compatibilized polymers are chemically bounded in a way that creates a controlled and stable morphology with a unified thermodynamic profile. In view of multiplicity andcontradictory requirements of various properties for most of the applications, almost all the commercial PBAs are made of two or more thermoplastics,elastomeric modifiers along with a series of compatibilizerswithmodifierscompounded together. A considerable number of blends have been appearing in the market regularly, some of which are listed in Table 9. Generally, in each of the PBAs a specific characteristic is being imparted by the dispersed phase and the resultant material has certain properties that are an optimized combinationof both the components. Some of the resultant features of the PBAs and the contribution of each phase is shown in Table 10.

Second Stage Table 9 List of the Engineering PBAs At the downstream of the extrusion process a suitable reaction catalyst, reactive diluent, e.g., crosslinking Manufacturers name Trade Materials monomer, is fed into the molten polymer mix. Bayblend Bayer PC-ABS Pulse Dow Chemicals Third Stage Cycloy GE Plastics Fabrication of end product in the moulding press. Trim-2O00 Monsanto Aruloy ARCO Geloy-xp 4001 GE Plastics PC-ASA Bayer Makroblend PC-PBT X. INDUSTRIALPERSPECTIVES Dow Chemicals Sabre Eastman Ektar The UnitedStates, Western Europe, and Japan consume Performance almost 88% of the performance plastics of the total world GE Plastics Xenoy production. Speciality plastics, such as PEEK, LCP, GE Plastics Noryl GTX Nylon-PPE polyimides, conductivity plastics, and polyetherimides, BASF Ultranyl etc., are exclusively consumed by them. In the era of Allied Signal Dimension the globalization of business, global players have little Ashahi MPPE Monsanto motivation to give up their dominant positionin the field Trim-IO00 Nylon-ABS Borg-Werner Elemid of engineering plastics, thus restricting the-flow of the Atochem Orgalloy Nylon-PP technology. Further technological trends are, therefore, Dexlon D & S Plastics likely to be driven toward the knowledge-based materiEnphite Chisso als rather than capital-intensive large-volume materials. DSM engineering Nylon-Polyolefins Akulon Speciality engineered plastics are, therefore, the frontDupont Selar runner of plastic materials for the next decade. The Dupont Nylon-Polyarylate Bexloy major challengesat the present state of the engineering D & S Plastics Dexcarb Nylon-PC plastics industrymay be identifiedas building up of own Duralloy Celanese PBT-POM technological base for negotiating on equal footing with GE Plastics Cycolin ABS-PBT technology owners, encouraging the development of Dylark ARCO SMA-PBT speciality plastics and their applications to encase intrinFirestone Steron sic technical strength, establishingjoint research and deK-Resin K R 0 3 Phillips velopment (R&D) ventures with the engineeringplastics AMOCO Mindel A & B Polysulfone-ABS manufacturers for exploiting biggermarkets, and developing a culture of cooperative teamwork in order to ASA = acrylate-costyrene-coacrylonitrile;POM = polyoxymethachieve the goals. ylene. For other abbreviations see text.

Polymer Blendsand Alloys

661

Table 10 Commercially Available En~neerinnPBAs: Contribution from Each Component

P o l y m e 'rr a d e lame

PBAs

-I

lylon/PPE 'lends

Polymer-I1

ILess Moisture Abs.

m I[Absorption II Low temperaturc Toughness

PC :enoy iE Plastics ClPBT rlloy

sensitivity Chemical Low Tern-

layblend PC Iobay ClABS dloy

Impact Resis

Balance of chemical, Resist. Low Temp.

-

Crystallinity Impact strength Rigidity,Shrinkage Stresscrack Resistance

Toughness J -

ABS1 Impact ASA Resistance Processibility Toughness Impact strength at Processing HDT BroadTemp.range Weatherability -Heat Resistance . Stiffness P

Sensitivity Chemical Resistance

rloryl PP0 ?E Plastics 'POIHIPS 4lloys

HDT Rigidity FlameResistance

Improved Processing Combination of Heat Resist.withdimenf sionalstability Toughnesswith

Processing Impact Resistance

HIPS

. I I

m

Advantages Disadvantages

The properties profile of some of the PBAs is illustrated in Table 11.

B. R&D Management of PBAs Development

Novel copolymerization and alloying technology were predicted to be the critical forces for shaping the devel-

opment ofnew materials in the 1990s. However, the reality of this technology indicates a period of cautious reassessment of the daunting technical and economic hurdles for the development of PBAs in view ofthe challenges of comDatibilization, cost performance profile in actial applications, and competition from recycling of single resins.

662

Kaith

al.

Table 11 Properties Profile of Some of the Engineering PBAs

Materials PC-ABS PC-PBT

Alloy Nylon-PP Nylon-PPE Nylon-ABS Properties gravitySpecific 1.09-1.12 1.06 1.17-1.22 1.04-1.13 1.08 Tensile-strength-psi 7200 6600-8000 6500-7600 6500-7500 8000 6800-7400 Elongation % 270-330 str. impactIzod 6.9-12 2.4-18 18 10 16-19 modulus Flex. x 10 E3 600 316330-370 250-300 250-300275 205-220135-190 HDT-"F160 0 2 6 4 psi 250 167 abs. Moisture 24 h 0.85 1.2

1.13

-

-

4

5 7.00

1.os

0.15

0.05

-

The technology required for PBAs is not less cornnical expertise required for development ofa successful plex than that neededto develop a totally new polymer. PBA virtually mirrorsthe commitments necessaryto inHowever, the cost-riskof developing alloys from existvent a new polymer with market application consideraing material stocks is still less than that for inventing tions. The following flow diagram (see below) briefly illustrates the practices and designs of PBAs, particunew polymers. R&DmanagementofPBAsdevelopment is anim-larlyofengineeringplastics. portant and complex taskas the time, money, and tech-

I Define Desired Component Properties .--) Selection

T

L

NO

T

I

Yes

No

Select Proper Distinctive Compare Properties vs. + Processing +.Stability of Technique Specification - Morphology I

I Economy Test 2

Check on mutual Compatibility (Compatibilization)

l

Desired Morphology Characteristics

I

Select Rheology of Blend

Strict compounding Manufacturing

Table 12 Polymeric Modifiers Used for Various Matrices in Engineering

Property Impact strength toughness Heat distortion temp. stiffness Chemical resistance Barrier properties Processibility PET

=

Modifiers

Matrix PVC, PP, PE PC, PA, PPE, PET NBR, MBA, etc. PC, PA, ABS, SAN

ABS, ASA, SEBS, Block copolymer, SEBS, EPR, EPDM,

Amorphous polymers Polyolefins PVC, PP, PE, PS, etc.

Crystalline polymers PA, PVOH, PVDC Liquid crystal polymers

PET, PPE, Arylates, PC, PS0

polyethylene terephthalene; SEBS = styrene-ethylene butylene-styrene; PS0

=

polysulfone. For other abbreviations see text.

663


PBAs are designed explicitly to meet the needs of specific applicationson the basis of their property-processing-cost performances. One polymeris incorporated into the matrix of other polymers to impart specificcharacteristics as per the requirement along withthe appropriate compatibilizer to ensure stress transfer in between phases. The polymer blend constituents and composition must, therefore, be selected on the basis of the compensation of properties, considering the advantages and disadvantages associated with each phase. Table 12indicates some of the components used as modifiers.

XI. APPLICATIONS OF PBAs PBAs technology established itself in the late 1980s as a premier fieldfor the creation of new engineering materials and has promised virtually unlimited combinations of resins that could target specific market applications while taming the processing difficulties and prohibitive cost of exotic polymers. During the 1990s novel copolymerization and alloying technology came into the field for fabrication of the materials with desired properties. PBAs are used commercially in various fields, such as the aerospace industry [58,59,156] and the military as well as in industries and domestic articles, such as adhesives [60-831, development of new materials with improved properties, such as gas and moisture barrier, resistance to heat andcold, chemicals, corrosion, abrasion and moisture, etc. [84-121,156-1641.

9. E. Rofael and H.-A. May, Proceedings of the Seventeenth International Symposium on Particle BoardCom10.

11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23.

posite Materials, (T. M.Maloney, ed.) WashingtonState University, Pullman, Washington, pp. 283-295 (1983). D. B. Agarwal and L. J. Broutman, Analysis andPerformance of Fibre Composites, Wiley-Interscience, New York (1980). P. Beardmore, et al., Science, 208: 833 (1980). M. R. Piggott, Load-Bearing Composite Materials,Pergamon Press, New York (1980). S. W. Tsai and H. T. Hahn, Introduction to Composite Materials, Technomic, Westport, Connecticut (1980). J. Delmonte, Technology of Carbon and GraphiteFibre Composites, Van Nostrand Reinhold,New York (1981). L. H. Sperling,Interpenetrating Polymer Networks and Related Materials, Plenum Press, New York (1981). D. Hull, Introduction to CompositeMaterials, Cambridge University Press, Cambridge(1981). F. J. McGany, ECT, 13: 968 (1981). J. S. Hearsons, MPE, 153(1982). G. Lubin, Handbook of Composites, Van Nostrand Reinhold, New York (1982). L. H. Miner, MPE, 152(1982). H. L. Peterson, MPE, 152(1982). R.P. Sheldon, Composite Polymeric Materials, Elsevier, New York (1982). C. H. Baskedis, ABS Plastics, Reinhold, New York

(1964). 24. G. A. Morneau et al., ECT, 1 : 442 (1978). 25. J. M. Lantz, MPE, 6 (1982). 26. J. A. Manson and L. H. Sperling, Polymer Blends and Composites, Plenum Press, New York (1976). 27. D.R.Paul andSeymour Newman (eds.), Polymer Blends, Academic Press, New York (1978). 28. 0.Olabisi, ECT, 18: 443 (1982). 29 W. M. Saltman, EPST, 2: 678 (1965). 30. J. P. Kennedy and E. Tornquist(eds.), Polymer ChemisI

ACKNOWLEDGMENT The authors acknowledge witha profound sense of gratitude and obeisance the generous guidance andconstant encouragement extended by the esteemed and erudite principal, Dr. R.L. Chauhan, Regional Engineering College, Hamirpur, India, during the preparation of this chapter. REFERENCES 1. C.

2. 3. 4. 5. 6. 7. 8.

W.Bert and P. H. Francis, Trans. N . Y.Acad. Sci. Ser., 11, 36(7): 663 (1974). J. E. Gordon, The New Science of Strong Materials, 2nd. ed., Penguin Books Ltd. H. Worth, UK, An Excellent Historical Perspective on Structural Materials andFine on Stiffness, Strength and Toughness (1976). P. J. Fenelon, Plastics Eng., 37(a):37 (1981). R. L. Jalbert, Alloys in MPE, 127 (1982). H. F. Mark et al. (eds.), Composites, Encyclopedia of Polymer Science and Engineering, 3rd ed., 3 (1985). H. F. Mark et al. (eds.), Composites, Encyclopedia of Polymer Science and Engineering, 4th Ed. 4 (1986). F. W. Billmeyer, Jr., Text Book of Polymer Sciences, 3rd ed., A Wiley-Interscience Publication, John Wiley and Sons, New York, p. 470.

31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41.

try of Synthetic Elastomers, Wiley-Interscience, New York, Part I (1968) and Part I1 (1969). I. D. Rubin, Poly (l-Butene):Its Preparation and Properties, Gordon and Breach, New York (1968). R. G. Bauer, ECT, 8 : 608 (1979). L. J. Kuzma and W.J. Kelly, ECT, 8: 546 (1979). J. E. McGrath, ECT, 8: 446 (1979). J. A. Brydson, RubberChemistry,AppliedScience, London (1978). H. W. Robinson, EPST, 6 : 505 (1967). H. L. Hsieh et al., Chem. Tech., 11: 626 (1981). D. C. Blackley,Synthetic Rubbers:Their Chemistry and Technology, Elsevier, New York (1983). M. G. Arora et al., Industrial Polymers, Polymer Chemistry, 1st. ed., Anmol Publications Pvt.Ltd., New Delhi, p. 277. R. W.Lang, et al. Polym. Eng. Sci., 22: 982 (1982). J. C. Seferis and L. Nicolais, (eds.), The Role of the Polymeric Matrix inthe Processing andStructural PropertiesofCompositeMaterials, Plenum, New York

(1983). 42. J. P. Trotignonetal., PolymerComposites, 3 : 230 (1982). 43. L. J. Buckley et al., SAMPE, Q : 16 (1984). 44. T. Mathews, Polymer Mixing Technology, Elsevier,

North-Holland, Inc, New York, Applied Science Publishers, Barking, U.K. (1982). 45. L. Mascia, The Role of Additives in Plastics, Edward Arnold Publishers, Ltd., London (1974).

664 46. J. A. Brydson, Plastic Materials, 4th Ed., Butterworth Publishers, Ltd., London (1982). 47. D.V. Rosato, Handbook of Composites, (G. Dubin, ed.), Van Nostrand Reinhold Co., New York (1982). 48. K.M. Prewo, Frontiers in Materials Technology, (M. A. Myers, ed.), Elsevier, Amsterdam (1985). 49. W. K. Fischer, U.S. Patent 3806,558,3835,201 (1979). 50. H. L. Moms, J . Elastomer Plast., 6 : 1 (1974). 51. R. J. Crawford, Plastic Engineering, 2nd ed. (1987). 52. M. 0.W.Richardson, Polymer Engineering Composites, Applied Science Publications, London (1977). 53. F. Rodriguez, Principles of Polymer Systems, 2nd ed,, McGraw Hill, International Ed., p. 258 (1983). 54. D. Mull, An Introduction to Composite Materials,Cambridge Solid-state Science Series, Cambridge University Press, Cambridge, Massachusetts, 3 (1981). 55. J. M. G. Couri, Encyclopedia Polymer Science Eng. (J. I. Krotschwitz, Exe. ed.),2nd. Editing SupplementVolume, Wiley Guter Science Publication, JohnWiley and Sons, New York, pp. 455-480 (1989). 56. L. A. Utracki et al., Polymer Alloys, Blends and Ionomers: An overview in Multiphase Polymers: Blends and Ionomers (L. A. Utracki and R. A. Weiss, eds.), ACS Symposium Series No. 395, 1-35 (1989). 57. S. Y. Kienzle, Advances in Polymer Blends and Alloys Technology (Kohudic, ed.), Technomic Publishing Co. Inc., Lancaster, 1, 1-11 (1988). 58. B. C. Hoskin and A.A. Baker, Composite Materialsfor Aircrafr Structures, AIAA Publ. Inc., New York (1986). 59. S. M. Lee (ed.), International Encyclopedia of Composites, 1-4, VCH, New York (1990). 60. P. J. Corish, Science and Technology of Rubber (F. R. Erich, ed.), Academic Press, New York (1978). 61. B. S. Gesner, Encyclopedia of Polymer Science and Technology. 62. S. N. Angove, Rubber J . 149(3): 37 (1967). 63. D. C. Blackley andR. C. Charnock, J . Inst. Rubber Ind. 7: 60 (1973). 64. D. C. Blackley and R. C. Charnock,J . Inst. Rubber Ind. 7 : 113 (1973). 65. A. King, RAPRA Res. Bull., 127 (1964). 66. M. H. Waltersand D.N. Keyte, Trans.Inst.Rubber Ind., 38: 40 (1962). 67. L. Bohn, Rubber Chem. Tech., 41: 495(1968). 68. K. A. Burgess et al., Rubber Age, 97(9): 85 (1965). 69. D. J. Angier and W. F. Watson, J . Polym. Sci., 18: 129 (1955). 70. D. J. Angier and W. F. Watson, Trans.Inst.Rubber Ind., 33: 22 (1957). 71. W. M. Hess, et al.,Rubber Chem. Tech.,66: 329 (1993). 72. N. Tokita, Rubber Chem. Tech., 50: 292 (1977). 73. G. N. Avgeropoulos, et al. Rubber Chem. Tech., 49: 93 (1976). 74. W.M. Hess, et al., Rubber Chem. Tech., 47: 64 (1974). 75. L. Evans, et al., Rubber Age, 9 4 : 272 (1963). 76. W.M. Hess, et al., Rubber Chem. Tech.,50: 301 (1977). 77. W.M. Hess et al., Rubber Chem. Tech., 58: 350 (1985). 78. P. A. Marsh et al., Rubber Chem. Tech., 40: 359 (1967). 79. P. A. Marsh et al., Rubber Chem. Tech., 43: 400 (1970). 80. H. E. Trexler et al., Kautsch Gummi Kunsts., 40: 945 (1987). 81. A. Y. Coran and R. Patel, Rubber Chem. Tech., 54: 892 (1981). 82. A. Y. Coran and R. Patel, Rubber Chem. Tech., 55: 116 (1982).

Kaith et al. 83. D.K. Setua, et al., Kautsch Gummi Kunsts., 44: 821 (1991). 84. R. R. Goddard, Monitor, 31: 1 (1990). 85. F. S. hoste, Monitor, 31: l1 (1990). 86. W. J. Kores, Barrier Polymers and Structures: Overview in Barrier Polymers and Structures (W. J. Kores, ed.) ACS, Washington, 1 (1990). 87, S. A. Stlrn and S. Tohalaki, Fundamentals of Gas Di#iusion in Rubbery and Glassy Polymers, ACS, Washington, 22 (1990). 88. P. M. Subramaniam, Polymer Blends: Morphology and Solvent Barriers, ACS, Washington (1990). 89. R. Bell et al., AutomotivePolymers and Design, 13 (1991). 90. W. M. Hess, et al., Rubber Chem. Tech. 66: 329 (1993). 91. A. Y. Coran and R. Patel, Rubber Chem. Tech., 53: 141 (1980). 92. A.Y. Coran, Rubber Chem. Tech., 63: 599(1990). 93. A.Y. Coran and S. Lee, Rubber Chem. Tech., 65: 231 (1992). 94. T. R. Wolfe, Thermoplastic Elastomers (N. R. Legge et al., eds.) Hanser, New York, Chap.6 (1987). 95. C.P. Roder, Handbook of ThermoplasticElastomers (B. M. Walker, ed.) Van Nostrand-Reinhold, New York, Chap. 4. 96. L. H.Sperling, IPN and RelatedMaterials, Plenum Press, New York (1981). 97. E. T. McDonel et al., Polymer Blends, 11, (D. R. Paul et al., eds.) Academic Press, New York, Chap. 19. Ghosh et al., Polymer Science, Sivaram,Tata 98. P. McGraw Hill, p. 871. 99. L. M. Glanville and P. W. Milner, Rubber Plast. Age, 48: 1059 (1967). 100. J. F. Svetlik et al., Rubber Age, 96(4): 570 (1965). 101. S. M. Hirschfield, U.S. Patent, 3,281,289(1966). 102. C. W. Snow, U.S. Patent, 3,280,876(1966). 103. V. Abagi et al., Mater-Plast.(Buchanest), 6(2): 84 (1969). 104. Chemische Werke HulsA. G. French Patent 1,499,094 (1966). 105. W.C. Flanigan, Rubber Age, lOl(2):49 (1969). 106. R.N. Kienle et al., Rubber Chem. Tech.,44: 996 (1971). 107. A. Springer, RAPRA Transl., 1170 (1964). 108. M. S. Sutton, Rubber World, 149(5): 62 (1964). 109. K. Satake et al., J . Inst. Rubber Ind., 4 : 102 (1970). 110. M. S. Baneji et al., International Conference Polymer '94, India IV, BCN-29. 111. D. D.DunnomandH. K. de Decker, RubberWorld, 151(6): 108 (1965). 112. B. Topcik, Mater. Plast. Elasromeri, 35: 201 (1969). 113. S. A. Banks et al., Rubber World, 151(3): 62 (1964). 114. A. K. Bhowmik et al., Rubber Products Manufacturing Tech., Marcel Dekker Inc., 693. 115. M. S. Baneji et al., International Rubber Conference, China (1992). 116. M. S. Baneji et al., IRMRA, 16th Rubber Conference (1993). 117. X. Zhang and Y.Zhang, International Rubber Conference, India, 29 (1993). 118. M. S. Baneji et al., International Rubber Conference, India, 39 (1993). 119. M. S. Banerji et al., International Rubber Conference, USA (1993). 120. C. K. Das et al., IRMRA, 15th Rubber Conference,India (1990).


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44 Reactive Compatibilization of Immiscible Polymer Blends Snooppy George, Josephine George, and Sabu Thomas Mahatma Gandhi University, Kerala, India

1.

on two demanding structural parameters [2]: a proper interfacial tension leading to a phase size small enoughto allow the material to be considered as macroscopically “homogeneous,” and an interphase “adhesion” strong enough to assimilatestresses and strains without disruption of the established morphology. Various techniques have been reported to combat the problem of incompatibility [3]. One way is to organize the phases and to increase their interfacial areas so that they are able to effectively transfer applied forces despite having low natural interfacial adhesion. Techniques of blending [4] and blending methods combined with certain block copolymersto stabilize the blend morphology [5] have been used successfully to form interpenetratingco-continuous networks of phases. Such systems have improved mechanical properties because the continuous nature of the phases allow direct load sharing of the components without the need for shear force transfer across phase boundaries. It iswidely known that the presence of certain polymeric species, usually suitably chosen block or graft copolymers, can alleviate to some degree these problems as a result of their interfacial activity [6-131. The segments of these copolymers can be chemically identical those with in the respective phases [8,13-151 or miscible withor adhered to one of the phases [16-181. As pointed out by Paul [l91 this type of surface activity should: (1) reduce the interfacial energy between the phases, (2) permit finer (3) provide a measure of stabildispersion during mixing, ity againstgross segregation, and (4) result in improved interfacial adhesion [20]. There have been many experimental investigations that studied the interfacial and surfactant properties of block copolymers. In additionto documenting the com-

INTRODUCTION

Blending of polymers has resulted in the development of polymeric materials with desirable combinations of properties. In mostcases, simple blending willnot result in the attainment of desirable combinations of properties because of some inherent problems. Frequently, the two polymers are thermodynamicallyimmiscibleand the product will not exhibit homogeneity. Oftena low level of inhomogeneity is preferred as both the components can retaintheir identity in the blend andthereby contribute synergistically toward blend properties. However, the interfacial situation in an immiscible blend is very critical in determining the blend properties. A typical incompatible blendexhibits high interfacial tension and poor adhesion betweenthe phases. The high interfacial tension results in poor dispersion during mixing andthe subsequent lack of stability during later processing or use. Often such blends show poor mechanical properties, particularly those related to ductility, that can preclude their commercial utilization.Moreover, the weak interfacial interaction in these blends leads to premature failure under stress as a result of the usual crack opening mechanism [l].

II. COMPATlBlLlZATlONCONCEPTS The two genericterms found in the blend literature are “compatibility” and “miscibility.”Componentsthat resist gross phase segregation and/or give desirable blend properties are frequently said to have a degree of compatibility even though in a thermodynamic sense they are not miscible. In the case of immiscible systems, the overall physicomechanical behavior depends critically 667

668

George et al.

patibilizing activity in polymer blend dispersion, these investigations havealso provided evidence of the interfacial activity in polymericsystems [21-251, surface activity [26-291, and dispersed efficiency [30,31]of block copolymers. For example, Gailard and coworkers [21,241 demonstrated the surface activity of block copolymers by studying the interfacial tension reduction in demixed polymer solutions. Addition ofpoly(styrene-block-butadiene)block copolymer to the polystyrene-polybutadiene-styrene ternary system first showed a characteristic decrease in interfacial tension followed bya leveling off. The leveling off is indicativeof saturation of the interface by the solubilizing agent.

a 2 0

II

I

CH-C-OH

l

II

I

CH- C-OH

REACTIVE COMPATlBlLlZATlON

Addition of block copolymers, crosslinking of rubber phase, or the generation of interpenetrating networks, which are being used for controlling morphology and mechanical properties of incompatible polymer blends, are not alwayspreferred. Therefore, more attention has been focused on reactive polymer processing in which superior polymer alloys are produced during melt processing,making it an attractive cost effective alternative. The basic principle underlyingreactive compatibilization isthat, by making use of the functionalitiespresent in one or more polymers, one can form graftor block copolymers in situ during processing. These copolymers act as compatibilizers by reducing the interfacial tension and increasing the adhesion between the phases and thereby allowing finer dispersion and more stable morphology. Several review articles [32-341 have appeared in the literature highlighting the significance ofreactive polymer processing. A.Polyolefin-PolyamideBlends Blends based on polyolefins have been compatibilized by reactive extrusion where functionalized polyolefins are used to form copolymersthat bridge the phases. Maleic anhydride modifiedpolyolefinsand acrylic acid modified polyolefins are the commonly used modified polymers usedas the compatibilizer in polyolefin-polyamide systems. The chemical reaction involved in the formation of block copolymers by the reaction of the amine end group on nylon and anhydride groups or carboxylicgroupsonmodifiedpolyolefins is shownin Scheme 1. One of the earliest references on compatibilizing a nylon-6-polypropyleneblendusingmaleic anhydride grafted PP (PP-g-MAH) wasthe work of Ide and Hasegawa publishedin 1974 [35].In their study, the formation of a graft copolymer was confirmed by D X after solvent extraction of the PP component. Blends with PP-g-MAH

J n

Nylond

Acrylic acid grafted polyolefin

I

-H20

CH2 0

I

111.

L

CH2 0

I1

CH -- C-OH

I

,

0 Maleic anhydride grafted polyolefin

l

1

-OH n

Nylon-6

CH2-C-OH

It

0 Block copolymer

Scheme 1 The proposedacidlamineandanhydride/

amine reaction between carboxylic groupslanhydride groups on modified polyolefins and amine end groups on Nylon-6.

showed a minimum in melt flow rate versus composition curve (varying the amount of N6), whereas blends without PP-g-MAH first showed a monotonic increase. There are several Japanese publications [36-381 and patents maintaining similar compatibilization of a nylon6-PPblendusingPP-g-MAH. A few studies [39-421 have been reported on laminates of N6 and PP and, in

pact

Compatibilizationof Polymer Blends

669

behavior to the greater reactivity of maleic groups than acrylic groupsto the amine end groups of nylon to form block copolymers. In the PP-PA system, the DSC thermograms showed two peaks corresponding to nylon and PP. For the compatibilized system the crystallization peak of nylon remainsunaltered, while that of PP shifted toward a higher temperature in the case of PPacr, and for PPmal, the shift was to lower temperatures (Fig. 1). This may be due to the fact that PPacr was acting as the nucleating agent. The average crystallinity of the blend was also decreased by the incorporation of compatibilizer. The mechanical properties of these blends was improvedby the addition of PPmaland PPacr as compatibilizers (Table 1). Generally, chemical reaction between the components of a blend increases the viscosity. Thus, an increase in viscosity in the presence of a compatibilizer may indicate that there are strong interactions at the interface. In the PP-PA system, Valenza et al. [43] reported such an increase in viscosity in the presence of a compatibilizer due to the increased interaction at the interphase. Miettinen and coworkers [44] also reported such anincrease in viscosity of the PP-PA-6 system by the addition of SEBS-g-MA as a compatibilizer. Here also the reaction at the interface is in between anhydride and aminegroups. The authors have evaluated the effectiveness of the SEBS-g-MA in PP-PAsystem in terms of impact property, morphology, and dynamic mechanical properties. In these blends, the addition of SEBS-g-MA leads to a very interesting morphology. In the ternary blends, in which PA-PP ratios are 20:80,40:60, a welldispersed phase of combined SEBS-g-MA and polyamide isseen in the polypropylene matrix on TEM investigation. A clear change in morphology is observed in a PA-PP ratio of 80:20 with 10% SEBS-g-MA. In this case the morphology showsa fine dispersion of size of about 0.04 p m and coarse dispersion of 0.1-1 pm in the continuous polypropylene matrix. Miettinen et al. [M] interpretedthismorphology as follows. In the blends of PA-PP 60:40 with 10 wt% SEBS-g-MA,the PNSEBSg-MA agglomeratesandbegins to overlap when the amount of PA is increasedto 80%. Therefore, the phase inversion inside the agglomerates extends to the whole

t 0

3

Figure 1 DSC thermogramsofNylon-12-PPblends. Source: Ref. 43.

particular, on improvingthe adhesion of these laminates by the use of modified PP or other modified polar thermoplastics in the packaging and automotive industries. The morphology, rheology, and crystallization behavior of reactive compatibilizedblends of polyolefin-polyamide blends were widely investigated. Polypropylene-polyamide compatibilization through functionalizedpolypropylenesis one of the systems studied recently. Valenza et al. [431 investigated the effectiveness of maleic anhydride grafted PP (PPmal) and acrylic acid grafted PP (PPacr) as the compatibilizer in the PP-PA system. In these systems, the average diameter of the inclusion isdecreased by the addition of modifiedpolypropylenesandthisreductionismore pronouncedin the case of PPmal.They attributed this

Table 1 MechanicalProperties

Nylon-l2 PP at at

Stress 35 Elongation at yield (%) Stress break Elongation at 200 (%) Izod Source: Ref. 43.

Nylon-12-PP Nylon-12-PP-PPmal Nylon-12-PP-PPacr

25

36 18

40 400

55

-

30

37

95

22 60 35

30 15

-

26

150

110

80

65

670

George et al.

sample. This results in a morphology with a bimodel dispersion of SEBS-g-MA as the continuous phase and PA-6 as a coarse dispersion of 0.4-1 pm. This morphology leads to exceptional impact properties (Fig. 2). The rheology and morphology of polyacrylamide-polypropyleneblendscompatibilizedthrough MA-g-PP wasstudied byHaddout et al. [45].In the compatibilized blend, the viscosity shows two types of behavior, i.e., at low shear rates up to IO3 S" a positive deviation is observed. The viscosity of compatibilized blends is higher than the uncompatibilized ones. The compatibilized blend also shows a decrease in domain size of the dispersed phase. The higher values of viscosity of compatibilized blends compared with uncompatibiked ones are due to the reduction of the average size of the dispersed nodules and also due to the better adherence between the two phases in the presence of compatibilizer. The in situcompatibilization of LDPE-polyamide-l1 was carried out by Lambla and Seadan [46]. The graft copolymer was formed by the thermal decomposition of selected peroxidesin the melt. The reactivity could further be enhancedby the addition of monomers, such as undecenal or maleic anhydride, both of which lead to significant changes in morphology and macroscopic properties in comparison with blends containing no such monomers. The kinetics of the reactive compatibilization of nylon-6-PP byacrylic acid modifiedPP was investigated by Dagli et al. [47]. The compatibilization reaction in this system involved the reaction between the acid group of acrylic acid modified PP and the amine groupof nylon6. A typical intensive batch mixer torque ( 7 ) vs time (t) trace for a ternary blend showing an increase in mixing torque upon the addition ofPP-g-AA to a binary PP-NBR (85:7.5) blend is shown in Fig. 3. The kinetic

I

0.0

09

0.4

0.6

0.8

PA6 WEIGHT FRACTlON vs PP

.

1.o

Figure 2 Notched Izod impact strength of ternary PA-PP-SEBS-g-MAH blends with 0 wt%, 5 wt%, and 10 wt% SEBS-g-MAH. Source: Ref. 44.

4 lmax

t

-

S= knax-

t

Omax

leqt

T

t Figure 3 Torque vs time trace for the PP-N6-PP-g-AA ternary blend. Source: Ref. 47.

studies revealed that the compatibilization reaction follows a second order kinetics as shown in Fig. 4.Interestingly, acid-amine nylon polymerization reaction also followssecond order kinetics [48-491. Also, they have investigated the effect of process variables on the compatibilization reaction. It was observed that changes in the rotor speedaffected the reaction kineticssignificantly, whereas temperature changes in the 15-20°C range have noeffect. The experiments in the twin-screw extruder had revealed that the screw speed, presence of venting, andsequence of feeding hada noticeable effect on the reactive compatibilization process. Venting and

0.0

0.2

0.4

0.8

0.8

S2 Figure 4 Reaction kinetics plot showing the use ofa differential methodof rate determinationof PP-N6-PP-g-AA ternary blend. Source: Ref. 47.

671

Compatibilization of Polymer Blends

residence time hada more significant effect on the grafting reaction. Screw speed was important from the distribution of the compatibilizer point of view. However, feed rate and temperature were not found to be significant in the range tested. Ionomers are sometimes used to compatibilize N6 and polyolefins blends. Carbonyl groups fromthe ionomer were foundto interact with the amine groups of N6 by forming hydrogenbonds. Ionomer concentrations as low as 0.5 wt% were found to be effective. Willis and Favis [50] have carried out extensive studies to determine the emulsifying effect of an ionomer (polyethylenemethacrylate acid-isobutyl acrylate terpolymer) on the morphology of polyolefin-polyamide blends. The addition of the ionomer to polyolefin-polyamide blends was found to increase the adhesion between the matrix and dispersed phase. Much finer dispersion of the minor phase in the matrix polymer as well as a reduction in the polydispersity of the particle size were also observed. One important finding was that as little as approximately 0.5% ionomer by weight of the blend was sufficient to produce a maximum reduction of the dispersed phase size. Larger quantities of ionomer yielded aggregates of the dispersed polyamide, possibly due to the clustering phenomenonof ionomers. B.

PS-BasedBlends

Baker and Saleem [SI] have reported on the reactive compatibilization of oxazoline modified PS andcarboxylatedpolyethylene. The coupling reaction results in amide-ester linkagesat the time of melt mixing. A schematic representation of the reaction isshown in Scheme 2. The graft polymer formed during melt mixing results in good interfacial adhesion between the phases. The

OPS-g-CPE

Scheme 2 Reaction scheme for the formation copolymer betweenOPS and CPE.

of block

intermolecularreaction between the two polymers furis ther substantiated by the rise in mixing torque, FTIR, and DSC studies. The Lewis acid initiated meltreactions of polystyrene and EPDM rubber to obtain high impact polystyrene were reported by Mori et al. [52]. It was found that among the Lewis acids investigated, the NaCl AICL double salt has the best initiating effect. The greatest number of crosslinks and the smallest extent of degradation was observed in this reaction. The extent of coupling increases with increasing Lewis acid and rubber content. As a result of coupling, the solubility of the elastomer particles in the continuous PSmatrixincreases. Increased solubility results in increased adhesion and mechanical properties. Willis et al. [53] have reported on the morphology andimpact properties of polystyrene-maleicanhydride-bromobutyl rubber blends as a function of interfacial modification and melt processingconditions. It was found that dimethyl amino ethanol (DMAE) serves as a reactive compatibilizing agent for these blends. The number average andvolume average diameters as a function of DMAE concentration for the PS-bromobutyl rubber (80:20)blend is shown in Fig.5. It is evident that the size of the dispersed rubber phase and polydispersity decrease as the extent of interfacial modification increases. This gives rise to a corresponding decrease in the interparticle distance. The particle size tends to equilibrate towarda larger concentration of DMAE.The impact strength values obtainedfor 80: 20 PS-bromobutyl rubber blend as a function of DMAE concentration are shown in Fig. 6.A three-fold increase in the notched Izodimpact strength is observed over the range of DMAE concentrations studied. It is very interesting to note that by increasingthe concentration of DMAE from 2-5% (i.e., by saturating the interface), the impact strength is significantlyenhanced while the particle size, polydispersity, and interparticle distance remain almost constant. This illustrates the importance of saturating the interface with the reactive compatibilizing agents. Teh and Rudin[54] have reported on the compatibilization of the PS-PE blend through reactive processing in a twin-screw extruder. In fact, this study [54] was aimed at generating the polystyrene radicals at the interface so that the coupling reaction could occur at the interphase between the PS andPE. Here compatibilization was achieved throughreactive processing with styrene monomer, dicumyl peroxide, and TAIC in a twinscrew extruder by simultaneous melt blending, polymerization, grafting,andcoupling. The polymerization, graftingkoupling and degradation reactions possible in a mixture of PS, PE, and styrene monomer (ST) in the presence of peroxide and the coupling agentare shown in Scheme 3. From thermal, molecular weight, and FTIRstudies on the blends and isolated PE and PS fractions, it was deduced that a polyalloy of PE-TAIC-PS was formed.

672

George et al. 0

Particle Size (urn)

l 6

0

Ob 0

U

L1

2

4

I

6

n T

10

8

% DMAEConcentration

Figure 5 Dependence of particle size of dispersed bromobutyl rubber on the DMAE concentration. Source: Ref. 53.

0

1

2 3 4 % DMAE Concentration

5

6


673

Radical~~Ons: ROOR

2RO.

ROOR + TAIC

2ROH + .TAIC.

RO. + PE

ROH+PE.

-

RO. + PS

W

RO. + ST

the improvement in properties of PS-PE blends via the reactive extrusion process in the presence of dicumyl peroxide and triallylisocyanurate (TAIC). In that case, the distribution of the radical generator and coupling agent was likely to be uniform throughout the whole melt rather than concentrated at the interphase boundaries. The reaction between the oxazoline group and the carboxylic group hasalso been exploited in the case of rubber toughening of polystyrene using acrylonitrile butadiene rubber having carboxylic groups [56]. The concentration of the reactive oxazoline groups was varied by mixing polystyrene with a copolymer of styrene and vinyloxazoline. The reactive blends showed significant improvement in impact properties. The 5% OPS blends showed an increase of 73% in impact strength over the nonreactive blend and then decreased with increasing oxazoline content (Table 2). The important factors that affect the rubber toughening are: (1) interfacial adhesion, (2) nature of the matrix, (3)concentration of the rubber phase, and (4) shape and sizeof the rubber particles. In the PS-XNBR blend containing OPS, due to the reaction between oxazoline groups of OPS and carboxylic groups of XNBR, the interfacial adhesion increases and as a result, the minor rubber phase becomes more dispersed. The immiscible blend needs an optimum interfacial adhesionparticle and size for maximumimpact property. InPS-XNBR, a very smallconcentration of OPS provides this optimum interfacial adhesion and particle size. The interfacial adhesion beyond this point does not necessarily result in further toughening.

ROH+PS. ST.

Po&merization: ST. + nST

Gr@ng/Coqling: PE, +PE.), PS,.

___,

+Psy

PE, - PEy PS, - psy

PE. + PS.

PE-PS

2PS. + .TAIC.

PS-TAIC-PS

2PE. + .TAIC. PE. + .TAIC. + PS

-

PE-TAIC-PE

-

PE - TAIC PS

Scheme 3 Reactions in a mixture of PS, PE,and styrene monomer in the presence of a peroxide and the coupling agent.

C. PET- and PBT-Based Blends

The formation of the polyalloy results in improvement in the performance of the blends. This system is similar to the production of high-impact polystyrene (HIPS) where a rubber is dissolved styrene in monomer and then polymerized in the usual way. Even though the impact strength of the compatibilized PS-PE blend was higher than that of PS, it was much less than that of HIPS. In another study, Van Ballegooie and [55] have confirmed

Reactive compatibilizationof engineering thermoplastic PET with PP throughfunctionalization has beenreported by Xanthos et al. [57]. Acrylic acid modifiedPP was used for compatibilization. Additives such as magnesium acetate and p-toluene sulfonic acid were evaluated as the catalyst for the potential interchange or esterification reaction that could occur in the melt. The blend characterization through scanningelectron microscopy, IR spectroscopy, differential scanning calorimetry, and

Table 2 Impact Results of Unnotched Charpy Specimens of PS-OPS-XNBR Blends

load

80 75 70

Peak Impact XNBR OPS PS (W%) energy (wt%) (wt%) 0

20 20 20 201.21

Source: Ref. 56.

100

60

(J) 0 0

5 10

20

0.86 2.05 3.54 2.74

(NI 1900

2120 2380 2180 1690

ratioDuctile 0.22 0.29 0.41 0.38 0.23

674

mechanical properties revealed that functionalized PP promotes fine dispersed phase morphology, improves processability and mechanical properties, and modifies the crystallization behaviorof the polyester component. The presence of the additives that are used to catalyze interchange or esterification reactions does not, in general, improve anyfurther the blend morphology,properties, or its processability. Kanai et al. [58] recently reported the impact modification of variousengineering thermoplastics (PPS, POM, and PBT). Functionalized elastomers were used as impact modifiers. They have correlated the impact strength withthe particle size andinterparticle distance of the modifiers. In the case of PBT with maleic anhydride functionalized ethylene olefin rubber (EOR), the plot of impact strength vs the particle diameter did not have any significant relationship, however, the plotting of impact strength vs interparticle distance gave a good correlation (Fig. 7).A critical interparticle distance approximately of 0.4 pm wasobtained for PBT-EOR blends.

George et al.

l o o ,

80-

60-

40

t ala

D. EPDM-BasedBlends

lntorprriiclr Dirtrnco t [ p m ] Amine-anhydride reaction has also been exploitedin the impact modificationof polyamide with maleic anhydride grafted EPDM rubber and with ethylene-ethylacrylateFigure 7 Notchedimpactstrengthvs.interparticledismaleic anhydride terpolymer [59,60]. Greco et al. [61] tance for PBT-maleicanhydridegraftedEORblends. have reported on the degree of grafting of functionalized Source: Ref. 58. EPR on mechanicalproperties of rubber modified PA6. The reaction between succinic anhydride grafted EPR and PA6resulted in the formation of a graft PA6-g-(EPRg-SA)copolymer. The degree of graftingwasvaried from 0.6-4.5 wt%. The results reveal that the morphology and impact properties were better when the degree 50 of grafting was higher. The Charpy resilience (R) as a function of the test temperature for PA6homopolymerand for ternary 40 blends at an increasing degree of grafting is shown in Fig. 8. From the figure, it isclear that pure PA6 exhibits 67 30 very brittle behavior with very low resilience values, ‘E which remain unchanged over the whole investigated 7 S temperature range. This type of behavior is further suba 20 stantiated from morphological observations. However, for blends of the same compositions but containing EPRg-SA as the rubbery phase, a large enhancement of imIO pact properties is observed with respect to pure PA6. From the figure, it is clear that R increases with the increasing degreeof grafting of the EPR-g-SA. FurtherJ more, all these blends show a marked variation in R -50 -30 -IO IO 30 r (‘C) values over the temperature range in which the behavior of the material changes froma brittle to a ductile mode of failure. The location of this transition temperature is Figure 8 Charpy resilience (R)as a function of test tema function of degree of grafting and blend composition. perature for PA-6homopolymerandPA-6-EPRbinary blends with increasing degree of grafting (DG), 0-0%, 0A shift toward lower temperature is observed with the 0.6%, A-2.4%, and “ 4 . 5 % DG. Source: Ref. 61. increasing degree of grafting value of EPR-g-SA.

c

n

b

675

Compatibilizationof Polymer Blends

The morphology development during reactive and nonreactive blendingof ethylene propylene rubber with nylon and polystyrene was investigated by Scott et al. [62]. For the reactive blends, MA-EPDM is used in nylon-EPM and oxazoline-g-PS, and maleic anhydride-gEPM are used in PS-EPM blends. The reduction in domain size with the time of mixing was investigated. They reported that the major reductionin domain size occurs at short mixing times. Inthe case of the reactive system the volume average particle diameter of the dispersed phase is reduced from -4 pm to 1 pm within the first 90 S of mixing time. For the reactive blend systems the mixing torque and temperature are higher than that of nonreactive systems. This is due to the chemical reaction occurringat the interface. In PA-EPM-MA, the reaction is in between the amine end groups of PA and anhydride groups on EPM-MA, and in PS-OX-EPMMA the reaction is in between the oxazoline groups of PS and anhydride groups ofEPM-MA. The chemical reaction between the polymers increases the molecular weight ofthe polymer, and, as a result, the mixing torque is increased. The increase in torque and temperature is less in PS-OX-EPM-MA compared to PA-EPM-MA, since the oxazoline-anhydridereaction is slow compared toamine-anhydride reaction. The morphology of the blends also shows a reduction in domain size for reactive blends thannonreactive blends. The reduction in size is also more pronounced inthe PA-EPM-MA system due to the more reactivity of this system. The possibility of improvingthe mechanical and impact properties of poly(~(-)-3-hydroxybutyrate)(PHB), which is highly prone to brittle failure, by melt mixing with functionalizedrubber has been investigated by Abbate et al. [63]. It has been found that the best results were obtained when succinic anhydridegrafted EPR was used. Here the chemical interactions taking place between the two components duringthe blending process result in good adhesion andthereby improved mechanical properties.

E. NaturalRubber-BasedBlends The reactive extrusion of polypropylene-natural rubber blends in the presence of a peroxide (1,3-bis( t-butyl peroxy benzene) and a coagent (trimethylolpropane triacrylate) was reported byYoon et al. [M]. The effect of the concentration of the peroxide and the coagent was evaluated in terms of thermal, morphological, melt, and mechanical properties. The low shear viscosity of the blends increased with the increase in peroxide content initially, and beyond 0.02 phr the viscosity decreased with peroxide content (Fig. 9). The melt viscosity increased with coagentconcentration at a fixed peroxide content. The morphology of the samples indicateda decrease in domain size of the dispersed NR phase with a lower content of the peroxide, while at a higher content the domain size increases. The reduction in domain size

lo'

f i

0 0 . 0 2 phr

IO'

S

0 . 0 3 phr

2 0.05 phr

IOf 10"

LO0

10'

IO'

l 0'

o(rad/sec) Figure 9 Complex viscosity of PP-NR blends vs. peroxide content of PP-NR blends at 220°C. Source: Ref. 6 4 .

has been correlated with the compatibilizing effect of the interpolymers,which were formed in situ during melt extrusion, and to the increased viscosity of the PP matrix bycrosslinking. The mechanical properties, such as yield strength and impact strength, also increase up to 0.03 phr peroxide content and then decrease. The extrusion of PP and polyisoprene in the presence of maleic anhydride was reported as a compatibilization method by Els and McGill[65]. They isolated the block copolymer formed during mixing using the solvent extraction technique and characterized by means of TGA, DSc, and IRspectra. The IR spectra of the xylene extracted fraction at 100°C showed the presence of peaks characteristic of PP, polyisoprene, and maleic anhydride. The DSC scan also had the evidence for the formation of a graft copolymer. The WAXS investigations show the formation of the &crystalline form of polypropylene along withthe a-form for the blends extruded in the presence ofMA. However, the addition of MA had little effect onthe percentage crystallinity in the PP phase. The impact properties of the blends were also found to improve by extrusion in the presence of MA. F. SEBSRubber-BasedBlends

The influence of maleic anhydride modified styrene(ethylene-co-butylene)-styrene(SEBS) triblock copolymer as a reactive compatibilizer in a nylon-6-SEBS blend was investigated by Wu et al. [66]. Whenthe maleated SEBS was incorporated into the PA-6-SEBS blend,

676

George et al.

the domain size of the dispersed phase decreased and became more uniform. The xylene extracted samples of the uncompatibilized blend showed two large bare cavities on SEM investigation. With the addition of 2 wt% of maleated SEBS in the blends, the morphology (N/cm2) showed holes with a peculiar shape, volcano-like and protruding fromthe plane of the continuous phase. The presence of the protruding holesreveals the presence of better interfacial interactions. The protrusion of holes arises from the release of SEBS chains of the discrete phase from entanglements formed fromthe reaction of maleated SEBS chains with the continuous polyamide phase. The impact strength of the blend wasalso significantly improved on the addition of MA-g-SEBS (Table 3).

Table 4 Adhesionstrength of

the LDPE and PDMS Rubber Blends Containing EMA Sample Adhesion strength codea PEo PE2 PE4 PE6 PEIO Si0

Si4 Si6 Silo

2.6 3.2 4.1 4.3 4.5 2.6 4.2 7.0 7.3

Thesubscriptsrepresenttheproportion of EMA in each phase. Source: Ref. 67.

G. Dimethyl SiloxaneRubber (PDMS)-Based Blends

Santraet al. [671investigated the effect of ethylene methylacrylate copolymeras a chemical compatibilizerin the low-densitypolyethylene-polydimethyl siloxane rubber blends. The IR investigationof the individual blend components and a blend containing 6 wt% EMA showed evidence for the formation of a graft copolymerbetween EMA and PDMS rubber. They had suggested a mechanism for the formation of the graft copolymer. The lap shear adhesion strength of the blends using EMA in one of the phases shows substantial improvement with an increase in the concentration ofEMA (Table 4). The tensile impact strength of the blends also improved significantly with the incorporation of EMA up to 6 wt%, and after that the increase is marginal.

The SEM investigationshows that the particle size of the dispersed domain size decreased from 3.3 to 1.1 pm with the incorporation of 6 wt%, EMA, and this indicates the increased surface area of the dispersed phase morphology. The increase in surface area led to effective compatibilization and responsible is for the increased adhesionstrength and tensile impactstrength of compatibilized blends.

H. NBR-BasedBlends Most of the commercially availablereactive compatibilized systems contain acidic functional groups. Reactive

Table 3 .me Formulation and the Impact Strength of the PA-6-SEBS-g-MA Blends

Recipe MA-g-SEBS SEBS PA6 (gMA/gSEBS)(wt%) (W%)

12

13 14 13

12

100 85 85 85 85 85 85 85 85 85 85 85 85 85

strength Impact MA ratiograft (wt%) 15

-

11

2 3 4

-

Source: Ref. 66.

(J/m)

14

1

11

11

2 3 4

14 13

154.6

1

2 3 4

12

-

1

1101.5 149.7 202.2

0.09 0.09 0.09 0.09 0.1 1 0.11 0.11 0.11 0.14 0.14 0.14 0.14

42.6 k 3.14 103.1 f 16.40 112.0 k 9.41 k 15.30 1157.8 k 75.20 1100.6 f 50.90 117.6 & 22.50 202.6 ? 33.30 -c 71.60 804.6 k 85.40 107.7 k 12.90 k 14.60 424.1 f 56.50 k 21.10

677


compatibilization of immiscible polymer blends via the incorporation of basic functional groups havealso been developed. A well-known example is the reactive PS developed by Dow Chemical Company [68]. Basic functional groupssuch as glycidyl methacrylate (GMA) have been grafted onto LDPE [69] and PP [70]. Recently, a novel process of making primary amine functionalized polyolefins by copolymerization of masked functional monomers has been patented [71,721. Recently, Baker et al. [73] have reported on the effectiveness of different basic functional groups for the reactive compatibilization of polymer blends. In this study, glycidyl methacrylate (GMA), 2-hydroxy ethyl methacrylate (HEMA), 2-hydroxy propyl methacrylate (HPMA), t-butylaminoethyl methacrylate (TBAEMA), dimethylaminoethyl methacrylate (DMAEMA), and 2isopropenyl-2-oxazoline (IPO) were melt grafted onto PP homopolymer. The effectiveness of these functionalized PPS as compatibilizers for PP-acrylonitrile cobutadiene-co-acrylicacid rubber (NBR)blendswas evaluated in termsof impact properties and blend morphology. The effect of the functionalized PPSon the impact energy of the PP-NBR blends is shown in Fig. 10. There were minor improvements in impact energy in the blends when HEMA and TBAEMA grafted PPS were introduced into the matrix phase. These improvements

may be due to possible hydrogen bonding interactions between either the hydroxy groupof HEMA or the secondary amine group of TBAEMA with the carboxylic acid group in the NBR rubber phase. The impact energies of PP-NBR blends having HPMA and DMAEMA grafted PPS in the matrix phasewere actually lowerthan blends without any functionalized PPS. This is due to the fact that the tertiary amine functionalitycannot form an amide linkage by reacting with carboxylic acid and the secondary hydroxy group in HPMA is much less active than a primary hydroxy group in esterification reactions. However, a substantial improvement in impact energies of PP-NBR blends withGMA or IPO functionalizedPPS is observed. The PP-NBRblendswent through a brittle-ductile transition as the concentration of the functionalized PPS inthe matrix phase reached a leveling off at 13 wt% inthe case of IPO functionalized PP and 25 wt% in the case of GMA functionalized PP. Up to a ten-fold improvementin impact energy was observed when the brittle-ductile transition was reached. A typicalforce-displacementimpact curve obtained using the instrumental impact tester for PP and PP-NBR blends withvarious amounts of GMA grafted PP in the PP matrix is shown in Fig. 11. The maximum loads for both PP and the nonreactive PP-NBR blend and their total displacement were small. The force decreased rapidly after the maximum force was reached, indicating brittle failure. As GMA grafted PP (13 wt%) was introduced into the PP matrix phase the maximum load and the total displacement were significantly improved. However, the force dropped quickly to zero after the maximum force was reached, again indicating brittle failure. Only when 25 wt% of GMA grafted PP

i

20

1 0

6

l2

18

-1

24

30

Pwccntage of Functionalized PP inMatrix

Figure 10 Effect of different functionalized PPS on the

impact energy of PP-NBR blends, B-HPMA, 0-GMA,ATBAEMA, U-HEMA, 0-IPO, and A-DMAEMA. Source: Ref. 73.

0.0

110 1:5 20 Disphccmont ( c m )

2.5

3.0

Figure 11 Force displacement curves of a-PP, b-nonreactive PP-NBR blend, C-reactive blend containing 13 wt% GMA functionalized PP, and d-reactive blend containing 25 wt% GMA functionalized PP. Source: Ref. 73.

George et al.

678

Table 5 Effects of Various Phenolic Modified Polypropylenes on the Properties

of NBR-Polypropylene Compositions Containing Amine Terminated NBR ~

1

2

50

-

Recipe

Polypropylene Phenolic modified polypropylene NBR 45 45 45 45 ATBN SP-I045 0.5 SnC12-2H20 0.5 0.5 curative before time Mixing addition (min)

5 3

4

-

-

50(A)

45 5 5 0.5 5

5 5

5

5

5

~

-

) 50(C) 5

5

5 5

5

5

30

-

10.2

10.91.9

54

54

5

Properties

strength, Tensile UB (MPa)16.515.3 9.8 9.6 atStress 100% strain v100(MPS) modulus, Young’s E106 (MPa) 106 107 105 149 elongation, Ultimate EB (%) 390 79 36 set, Tension (%) break, stress at True vi$ (MPa) 173 7 5.1 5

20.1 420

.O

85.8 110.55

450 54

Source: Ref. 74.

was addedto the PP matrix phase didthe blend become ductile withthe force decreasing slowlyin a stable manner afterthe maximum force was achieved, whileat the same time the total displacement increased significantly. Coran and Pate1[74]investigated the reactive compatibilization of PP-NBR and HDPE-NBR blends using phenolic modified polyolefin, maleic anhydride modified polyolefin, and amine terminated nitrile rubber as reactive components. Dynamic vulcanization was also inves-

tigated as a way of reactive compatibilization. The mechanicalpropertiesandpercentageoilswellingwas of these improved by thereactivecompatibilization blends (Table 5). The improvement in mechanicalproperties and the decreasein oil swelling in reactive blends are attributed to the increaseininterfacialadhesion caused by the graft copolymer formed in situ during mixing. The graft copolymer is formed by the reaction between amine end groupsof NBR and the anhydride and

OH C H z e c b o H

+ N H R + E l AMINE TERMINATED NBR

PHENOLIC MODIFIED PP. 4

MALEIC ANHYDRIDE GRAFTED PP.

COMPATlBlLlZlNG BLOCK CO-POLYMER

Scheme 4 Reaction scheme for the formation of graft copolymer between amine terminated NBR and modified polypro-

pylenes.

inte~acialadhesion and to the red~ctionin inte~acial olic ~ o d i f i e dpolyolefin reacts with to form a block copolymer as shown efins are used as comoncentration of these

This block copolymer acts as an e ~ ~ l s i f y i nagent g in the blends leading to a red~ctio and im~rovedadhesion. At conc the critical value, eases the domain size

omain size o

of 70 :30 ~ c a ~ nelectron i n ~ micro~rap~s c anhydride modi~ed~olyethylene.Sourc

R blend with (a) 0 wt%, (b) 1 wt%, (c) 5 wt%, and (d) 10

680

George et al.

7

of compatibilizer below CMC, and a leveling off is expected after that. The expression for interfacial tension reduction (Av) in a binary blend upon the addition of divalent copolymer A-b-B is given by

?

Av

= d&( 1/2x

Ad

=

+

l/Zc) - 1/22 exp ( 2 4 2 )

(1) where d is the width at half height of the copolymer profile by Kuhn statistical segment length, & is the bulk volume fraction of the copolymer in the system, x is the Flory Huggin’s interaction parameter between A and B segments of the copolymer, and Zc is the degree of polymerization of the copolymer. Based on this equation, the plot of interfacial tension reduction versus & should yield a straight line. Since interfacial tension reduction is directly proportional to the particle size reduction, the interfacial tension reduction term in Eq. (1) can be replaced by the term particle size reduction. Therefore,

5 5

u W I-

w4

35

f3

$2 A

1

Kd& (1/2x

+

l/Zc)

-

1/Zc exp ( 2 4 2 ) (2)

0

0

5

10

15

WEIGHT % OF COMPATlBlLlZER

Figure 13 Average diameter of dispersed NBR domains vs. phenolic modified polypropylene and maleic anhydride modified polypropylene concentration. Source: Ref. 75.

where K is the proportionality constant. In HDPE-NBR and PP-NBR blend the particle size reduction (Ad)varies linearly with blend composition at low compatibilizer concentrations, whereas at higher concentrations it levels off, in agreement with the theories of Noolandi and Hong (Fig. 14).

fYCWH R

H3C

QUINONE METHIDE

DIMETHYLOL PHENOLIC COMPOUND

1

14

c

C-

) POLYPROPYLENE

R

QUINONE METHIDE

w- CH II

R

PHENOLIC MODIFIED PP

NITRILE RUBBER

Ph-PP

Scheme 5 Reaction scheme for the formation of graft copolymer in PP-NBR blends in the presence of phenolic modified

PP.


681

R OR

Scheme 5 Continued

The mechanical properties of these blends are also increased with an increase in compatibilizer concentration. The improvement in tensile strength and tear strength on the addition of modified polyolefins are shown in Fig. 15. The mechanical properties also show a leveling off after the optimum concentration. The im-

provement in strength with compatibilizer concentration is due to the improvement in interfacial adhesion, which leads to better stress transfer between the matrix and the dispersed phase. Ramesh and De [78] have investigated the effect of carboxylated nitrile rubber as a reactive compatibilizer for immiscible blends of poly(viny1 chloride) and epoxidized natural rubber. In PVC-ENR binary blends there is crosslinking due to interchain reactions. These blends are immiscible at various blend ratios. The incorporation of XNBR into these systems increased the compatibility of these blends. The IR in-

0 HAPE 4

A

PhPE

130

21

120

20 a

E E

%

1

z

E c9

I^

110

za

0

19

I-

cn

$100

W

9

cn

2

z

I-

0

-

0

A

1

1

I

1

2

4

6

8

1

10

10

W I-

90

A

I 12

b$

1

1

14

16

& VOLUME FRACTION OF COMPATIBILIZER

Figure 14 Effect of volume fraction of maleic modified PE and phenolic modified PE on the particle size reduction of 70 :30 HDPE-NBR blend. Source: Ref. 75.

80 0

5

10

17 15

WEIGHT % OF Ph - PP

Figure 15 Effect of phenolic modified PP concentration on tensile strength and tear strength of 70:30 PP-NBR blend. Source: Ref. 76.

682

George et al.

6). The dynamicmechanical analysis of the ternary blends with75 :25 PVC-ENR b l h d showed single Tg at all levels of XNBRconcentrations, which indicates the miscibility of the system. In the 50:50 PVC-ENR blend, when the concentration of XNBR increased, the blend becomes progressively miscible. IV. SUMMARY

I

CH3 H

I

I

........ CH-CHz... Cl

I

I

0

I

CH3 H

I

(B)

I (B) (C) ..'...CH2XHXH 1 , the opposite is true. Incompatible systems, such as most TP-TLCP blends, belong to the first kind of flow behavior. But it was also reported that some TP-TLCP blends havecomposition-dependentviscosityindicatingtwo minima below and above the 50 wt% ofTLCP. An example for this behavior wasa polysulfone-TLCP blend, its flow curves were measuredat 240°C [9]. In allthe cases, they must be able to describe their special flow behaviors by applying Eq. (7). The viscosity drop becomes more pronounced the in wholerange of 4 (0% < 4 < 100%) when K(l - A) increases, indicating that the value of parameter K considerably affects the flow behavior of the blend (Fig. 3).Note that K(l - A ) is related to the fibrillation and migration of TLCP-rich interlayer. As demonstrated, Eq. (7) gives complete information on howthe weight fraction influences the blend viscosity by taking into account the critical stress ratio A, the viscosity ratio 6, and a parameter K, which involves the influences of the phenomenologicalinterface slip factor a or ao,the interlayer number m,and the d/Ro ratio. It was also assumed in introducing this function that: (1) the TLCP phase is well dispersed, fibrillated, aligned, and just forms one interlayer; (2) there is no elastic effect; (3) there is no phase inversion of any kind; (4) A < 1.O; and (5) a steady-state capillary flowunder a constant pressure or a constant wall shear stress.

Yi

688

.

-K(I-X)=lO -K(l-X)=20

c=0.8; 6 = 1

... K(l-X)=30

1 eo

2

~

”””””c

’’;:.

.......,

0.1

0

20

40

60

80

100

TLCP Weight Fraction (%)

Figure 3 Theoretical viscosity function patterns by varying the parameter K(l - A).

C. ExperimentalResultsandDiscussions The TLCP used was KU 9231 produced by Bayer AG, Germany. The matrix material wasan engineering plastic polyethersulfone (PES) manufactured by Jilin University, China. Inan earlier article [1 l] we reported that KU 9231 was incompatible with the PES. The materials were dried in a vacuum oven at 115°C for 24 h. They were then melt blended by using a domestic twin-screw extruder (435) [screw diameter = 35 mm]. The weight ratios of PES-TLCP were 90: 10 and 70:30, respectively [12]. Figure 4 shows viscosity versus shear rate results for the two original components and their blends, re-

PES/TLCP

spectively. Rheologicalexperiments were conducted on an Instron Capillary Rheometer, Mode 3211. The data were systematically corrected according to the Bagley and Rabinowitschprocedures. The test temperature was 310°C. Prior to testing, the blended materials were dried again in the vacuum at 125°C for 4 h. As mentioned previously, andas expected, the mixtures containing a TLCP component exhibitedshear viscosities lower than those of their original components, whether the TLCP weight fraction was 10% or 30%. The higher the TLCP weight fraction, the lower the blend viscosity. The viscosity curves of the two pure components crossed each other at a shear rate about 25 s”, i.e., at this point the viscosity ratio was 1. The effect of the blending ratio on the melt viscosity of the two-phase system canbest be seen from the cross plots at different shear stress. For this purpose, the rheological data were first fitted to Eq. (7) by a computer programming, andthen, for T = q y being fixed, both the theoretical blend viscositycurves and the experimental values were plotted againstTLCP weight fraction in Fig. 5. It was foundthat the computer-fitted curves ran perfectly through the experimental data points. The calculated parameters K(l - A), E , and S are listed in Table 1, according to dependence on the shear stress. It is obviousthat all values of K ( 1 - A ) were found to be largerthan 30. They increased withincrease of the shear stress (i.e., shear rate), and ithas been well known that high shear stress leads to the formation of dispersed TLCP fibrils with a higher aspect ratio. Experimental evidence is given for the same PES-TLCP system in Yi et al. [7],where a larger average fibril aspect ratio was microphotographed at a higher shear rate. The average value of the exponent was approximatelyE = 1, indicating that the viscosity curves were nearly symmetric.

+100/0

4-90/10

+,70/30

-

a0/100

U U I

10’

I

Id

0

103

20

40

60

80

100

TLCP Weight Fraction (ss)

Shear Rate 7 (S”)

Figure 4 Experimental viscosity curves of PES, TLCP,

and their blends.

Figure 5 Comparison of theoretical viscosity curveswith experimental value points of PES/TLCP blends for fixed shear stressT.

689

Thermotropic Liquid Crystalline Polymer Table 1 Parameters Provided by Fitting the Experimental Data to Eq.(7) for PES-TLCP Blends

W‘a)

K(l - A)

80 100 125 200

2.259 43.43 45.94 55.78

x

E

6

1.03 1.14 0.90 0.88

1S83 0.962 0.838

The influenceof viscosity ratio S on blend viscosity reduction can be well described by Eq. (7).As an example of E < 1, it is seen from Fig. 6 that the theoretical viscosity curves have two ranges, both are completely below the order of unity. In the range of 6 < 1, particularly when 0.1, there is a considerable increasing tendency of 7/70against S, suggesting that the lower viscosity component (TLCP) may migrate to the capillary wall, controlling the flow behavior of the blends. The higher the shear stress and rate, the stronger this viscosity reduction effect. At 6 ratios greater than 1, i.e., at lower shear stress, the matrix becomes the lower viscosity component and may begin to control the blend flow. Thus, the 7/q0ratio will be less affected by S. Our investigation was performedjust near the turning rangeof 6. Equation (7) depicts the viscosity decrease independent of the chemical features of materials. Alsofor fixed 7 , Figs. 7 and 8 demonstrate a further example of a polyamide-TLCPblendwithdifferentweightratios. The rheological data in Fig. 7 were taken from Siegmannet al. [l]. It is obvious that the lowest blend viscosity is obtained at a TLCP loading of only 5%. This result is

.

, 10’

Id

,

103

Shear Rate -f

..J

104

(S”)

Figure 7 ExperimentalviscositycurvesofPA,TLCP, and their blends of different blending ratios[l].

somewhat different than the PES-TLCP blend previously reported. However, the weight fraction dependence of blend viscosityalso followed Eq. (7) quite well (Fig. 8). The parameters were also evaluated by fitting the rheological data to Eq. (7)and are listed in Table 2:It appears to be natural that both the E and S values were significantly lowerthan those of the PES-TLCP blend, previously discussed. It is interesting to note that our derivation of the viscosityfunction Eq. (7)wasnot restricted within TP-TLCP systems. The negative deviation flowbehav-

0

104 L

1

so 0.1 0.01

0.l

1

10

V% Figure 6 Theoretical viscosity function versus viscosity ratio S = qJq0 at different blending ratios 4.

0

I

0

20

40

‘60

80

100

TLCP Weight Fraction (%)

Figure 8 Comparisonof theoretical viscosity curveswith experimental values of PA-TLCP blends for fixed shear stress 7 .

Yi

690

Table 2 Parameters Provided by Fitting the Experimental Datato Eq. (7) for PA-TLCP Blends

.r(kPa) 300 340 387 440 500

K(l - A ) 29.00 1.384 32.90 38.39 0.814 43.73 0.391 47.14 0.391

6

E

0.03 0.06 0.10 0.14 0.15

1.100

ior fromthe log-additivity rulehas been also widely observedondifferent blends, for examplepolycarbonate-polyolefin blends reported by Huang et al. [ 141. They foundthat in both PC-HDPE and PC-PP systems the addition of 5% polyolefins to 95% PC considerably reduces the blend viscosity to a value lower than that of both parent polymers. Utraki and Sammut[l51 examined the morphology of a LLDPE [linear LDPE] blend containing 25% PC and reported that the low-viscosity component migrates to the capillary wall and controls the flow behavior. As a result, they also observed the viscosity reduction. In any case, an interfacial lubrication or an interfacial slippageof these incompatible systems is oftena reasonable concept for understanding the viscosity reduction. It is hoped that the function in Eq. (7) can serve as a possible basic tool to deal with the viscosity reduction of the two-phase systems. 111.

FIBRILLATIONBEHAVIOR AND RHEOLOGICAL DETERMINATION

Introduction A. Theoretically, the fibrillation behavior of a dispersed phase in a matrix isdictated by many characteristic fac-

tors such as Webernumber W,, viscosity ratio 6, breakup time t b , first normal stress difference, and the flow types. Under certain conditions [16-181, when discrete domains of a well-dispersed minor phase are deformed, elongatedellipsoidaldomainsmayform to nearly cylindrical domains.These domains, in turn, disintegrate into smaller drops, if a breakup time of tb is exceeded. So, the disintegration andcoalescence determine the morphology. Furthermore, a necessary condition for formation of fiberlike structures in the non-Newtonian range isthat the first normal stress difference of the matrix phase should be larger than that of the minor phase at the same shear stress [19]. According to this model, the deformation behavior of fiberlike drops in simple shear flow has been demonstrated with many, mostly incompatible, blends [203. In this section, we examine fibrillationbehavior of a polycarbonate (PC)-TLCP blend by injection molding. B.TheoreticalAnalysis Concerning a liquid droplet deformationand drop breakup in a two-phase model flow, in particular the Newtonian drop development in Newtonian median, results of most investigations [16,21,221 may be generalized ina plot of the Weber numberW , against the viscosity ratio 6 (Fig. 9). For a simple shear flow (rotational shear flow), a U-shaped curve with a minimum corresponding to 6 = 1 is found, and for an uniaxial extentional flow (irrotational shear flow), a slightly decreased curve below the U-shaped curve appears. In the following text, the U-shaped curve will be called the Taylorlimit [ 161. Rumscheidt and Mason [IO] studied the drop fracture of Newtonian fluids accordingto the value of a dimensionless 6. They found that 0.1 C 6 < 1.0 is most favorable for extensional deformation andsubsequently

1000

We

at Burst,

dimensionless

100

10

Irrotational shea! (4-roll)

1 \ “ ” ”

J

0.1

10-7

10-6

10-5

10”

loo

IO‘

lo2

103

Viscosity Ratio, &=‘ID/ ‘IC, dimensionless Figure 9 W, - 6 plot: comparison of effect of viscosity ratio 6 on critical shear We,criin rotational and irrotational shear

fields [18].

691

Thermotropic Liquid Crystalline Polymer

breakup. The exact W, - 6 graphs vary somewhat from author to author, but the general criteria is the same. Intermediateflow types between the two "regular" flow types give critical Weber numbers in between the curves

1211.

Table 3 Main Properties of Experimental Materials

TLCP P (&m3) E (MPa)

Both dimensionless Weber number and viscosity ratio are defined by:

PC (Sinvet 303)

13000

W " ) 156 168 HDT/A ("C) Elongation at break (%)

Ub

where 710 and V C stand for the viscosity of the dispersed and the continuous phase, respectively, in our case T~ = 71 and TC = 770. Representing the shear rate, shear stress, droplet diameter, and interfacial tension coefficient are y , Q , d. and U, respectively. Based ona lot of experimentalobservations, criteria for the drop stability can be defined as below the U curve, namely W, < We,cri,the interfacial stress can equilibrate the shear stress, and the drop will only deform intoa stable prolate ellipsoid. Abovethis curve, the viscous shear stress becomes larger than the interfacial stress. The drop is at first extended and finally breaks up into smaller droplets. The droplet deformation is also time-dependent. A practicable dimensionless breakup time t g can be defined as:

Grace [l81 evaluated tB as a function of 6 for viscoelastic matrices. His results are shown in Fig. 10. The reduced time increases as the viscosity ratio increases. If the residence time during a processing is less than

(Vectra A 950)

1.20 2500 65 130 >l50

1.40

2.6

the breakup time, for the irrotational shear or rotational shear, the elongated drops will keep their form. C. Experimental

The TLCP used was a widely investigated one, Vectra A950, manufactured by Hoechst-Celanese. The PC was produced by Enichemunder the tradename Sinvet 303. Table 3 lists the main properties of these materials according to the manufacturers. The materials were dried in a vacuum oven at 120°C for 4 h before processing. Prior to injection moldingof the composite samples, two components were melt mixed by a ZSK, W&P, Stuttgart, extruder with a constant weight ratio of PC-TLCP = 80:20 [23]. The standard injection molding machine used had a screw diameter of 30 mm and the aspect ratio of 23.70. The barrel temperature profile was 270, 280, 290, and 295°C. The mold temperature was about 90°C. The injection molded tensile samples were processed accordingto the CAMPUS specification (Computer Aided Materials Preselection by Uniform Standards) [24] and DIN53455 Form 3. To obtain the different flow conditions, four groups of samples were injection molded by varying melt

1000

Reduced Time

100

at W , , cri t:

10

dimensionless 1 L

0.1

1 0 ' ~ lo4

10"

10"

loo

IO'

lo2 lo3

Viscosity Ratio, S dimensionless Figure 10 f $ - S plot:comparison of effect of viscosityratio 6 onreducedbreakuptime rotational and irrotational shear[18].

t$

atcriticallimit

for

692

Yi Table 4 Processing Condition and Mechanical Properties of PC-TLCP Composite

Samples InjectionMolded properties Tensile condition Processing No.

Tm V3

1

280 320 280 320

104 2 3 95 81 4

Q (cm3/s)

Y( W

116 4648 116 4092 1158 3704 1158

112 8 8 80 80

E, (MPa)

ffb

(MPa)

4922

Mold temperature = 90°C; maximum clamping force = 700 kN;maximum injection pressure = 225 MPa.

temperature T,,, at the die and injection volume flux Q (Table 4). Hence, these samples will be identifiedin the following text by sample groups 1 , 2, 3, and 4, respectively. The average shear rate was calculated using the Eq.: 6Qe/(bh2),where Q is the volume flux, e = 0.772 [ 8 ] , b and h are the width ( b = 10 mm) and thickness ( h = 4 mm) of the cross-section of the dumbbell samples. Similar to prepared metallographicsamples, the injection molded samples were cut along the flow direction, smoothed, and polished in order to expose their internal surface. After proper etching, the treated surfaces of the flank cross-sectionwere photographed using a polarized light optical microscopy. Based the on color differences betweenthe TLCP and matrix, volume fraction and aspect ratio of the TLCP fibers were measured [W.

10

The mechanical properties were obtained using a tensile machineat room temperature and for a strain rate of lOOO%/h. Each reported value of the modulus wasan average of five tests. The tensile modulus E, was taken as the slope of the initial straight line portion of the stress-strain curve. D. Morphological andTensileResults

The results of viscosity versus shear rate are reported in Fig. 1 1 for the two pure components and their blend, respectively. The temperatures were the same for the viscosity measurements and for the injection molding. At temperatures of 280°C and 320°C, the viscosities of the blend are found to be values between the limits of the two pure components. In both cases, the TLCP still

100

1000

shear rate (l/s)

Figure 11 Viscosity curves of PC, TLCP, and their blends with respect to the test temperatures (* from WLF equation).


693

had the lowest viscosity. It was also observed that the pure TLCP exhibited a pronounced die shrinkagein contrast to the die-swell of most known thermoplastics. Figure 12 shows the dependence of the average aspect ratio and the TLCP volume fraction on the relative sample thickness for the four processing conditions in the core layer, transition layer and skin layer, respectively, by a morphological examination [13]. Generally, the aspect ratio increases from core to skin layer, whereas the situation is reversed for the volume fraction. An average volume fraction about 20% canbe clearly seen. Table 4 also reports dependence of the mechanical tensile properties of the samples on the processing conditions. The highest tensile properties of sample 1, injection molded with a lower melt temperature and a lower volume flux, are attributed to the highest degree of fibrillation of the TLCP fibers, as shown in Fig. 12, by socalled in situ reinforcement.

Processing Conditions

Calculating

Q, T,r, ...

Sample Geometry , ... Flow Curves

Calculating 6(x)

" Morphological Examination I/d(x)

H

Calculating I/d=D(6,We)

Calculating Evaluating We(&)

I

W&)

l Evaluating $((a)

Calculatingti(x)

Figure 13 Calculationprocedure.

Sample groupl

20

10

0 2

20

P F

10

B

o

0

E Cl

E. CalculationandDiscussion

20

10 group Sample

30

_.

30

20 10

g

The purpose of our calculation was to quantitatively evaluate the deformational behavior of the TLCP droplets and their fibrillation under the processing conditions, and finally, to establish a relationship among the calculated Weber number, the viscosity ratio, and the measured aspect ratio of the fibers. Figure 13 illustrates this procedure. All calculated results were plotted as

Y

S C

BV

E

L

""""_

" " " " "

/

c

0.4

0.6

kj (U c v)

500

-I

Halved sample thickness

0

0.2

0.8

1

normalized halved thickness of samples

Figure 12 Averageaspect ratio andvolumefractionof TLCP fibers as functionsof halved relative sample thickness for four processing conditions.

Figure 14 Calculated shear rate as function of the halved sample thickness for two injection volume fluxes Q.

Yi

694

functions of halved relative sample thickness with respect to the four processing conditions. At first, the theoretical shear rate curves fix) were calculated according to: ( m + 2) Q fix) =

(ir+'xm

quasi-steady flowstate, constant thermal physicalproperties, homogeneous melt, and isothermal flow conditions [8]. As shown in Fig. 14, the shear rate increases in both cases from central to outside, and its gradient strongly depends on the volume flux Q. (The dashed lines in the figure represent the average volume flux.) From the fix) functions and the two melt temperatures used, andby usingthe viscosity curves from rheological examinations (Fig. 1l), viscosity distributions ~ ( x of ) the two pure components were easily determined, as showninFigs.15aand15b. Subsequently, the viscosity ratio functions 6 ( x ) were also calculated (Fig. 16). All four curves fall slightly from the core to the outside.

(12)

bh

for two injection volume fluxes of Q = 80 cm3/s and 8 cm3/s, respectively. Where m is the flow index of the viscosity curves, and it wasan average value measured at 280°C and 320"C,and x is the thickness variable. This expression is onlyvalidwhenassuming a steady or

h

.e

100

g

s

'5

1

0 (4

0.2

0.4

normalized halved

0.6 thlckness of

0.8

l

samples

0 0

.v, 5 10

0 (b)

0.2

0.4

0.6

0.8

1

normalked halved thickness of samples

Figure 15 Calculated viscosityas function of the halved sample thicknessfor twomelt temperatures and for: (a) injection volume flux of 8 cm3/s; and(b) injection volume flux of 80 cm3/s.


695

10

0.01

0

0.2 0.4 0.6 0.8 normallzed halved thickness of samples

1

Figure 16 Calculated viscosity ratio as function of the halved sample thickness for the four sample groups.

To constitute the W, number, characteristic values such as the drop diameter, d, and particularly the interfacial tension, W , must be experimentally determined. However, the W , number can also be obtained by deductionfrommathematical analysis of droplet deformational properties assuming a realistic model of the system. For a shear flow that is still dominant in the case of injection molding, Cox [25] derived an expression that for Newtonian fluids at not too high deformation has been proven to be valid: D =

5(19S

4 (S

+ 16)

+ 1)[(196)* + (20W,)2]1'*= D(6y we)

in which, according to Taylor [16]: D = -l - d l + d where l and d are the long and short axis of a deformed droplet. This expression can be rewritten as follows

1

-=

d

+

1 D(S,W,) 1 - D(S,W,)

where lld is the aspect ratio that, in our study, had already been experimentally determined as lld = f ( x ) ,in Fig. 12. With a computer program, which solved Eqs. (5) and (7)for W, by numerical method, Weber numbers were obtained from layer to layer for each injection molding operation. The calculated points were then connected to curves, depending on the normalized thickness. Figure17 shows these theoretically deduced W,(x)

curves based on the known aspect ratio and the known viscosity ratio data. In general, the W , curves have an increasing tendency,'whereas for S(x),the opposite is true. To evaluate the fibrillation behavior of dispersed TLCP domains according to the W, - S relation discussed previously, different W, - S graphs were calculated by eliminating the thickness variable x. The result is reported in Fig. 18. It is obvious that all the points obtained are found to be relatively close to the critical curve by Taylor. The Taylor-limit is also shown in the figure with a solid curve. One finds that all the values calculated on sample 1 are completely above the limit, while all those determined on sample 4 are completely below the limit. The other two samples, 2 and 3, have the W, - S relation just over the limit. According to the criteria, the dispersed phase embedded in the matrix of sample 1 must have been deformed to a maximum aspect ratio and just began or have begunto break up. By observing the relative position of the experimental data to the critical curve, the deformational behaviorof the other samples can be easily evaluated. Concerning the fibrillation behavior of the PC-TLCP composite studied, the Taylor-Cox criteria seems to be valid. More general information of the fibrillation behavior in the form of average aspect ratio against both the viscosity ratio and the W, number is illustrated in Fig. 19 by using a mathematical smoothingoperation. Generally, in the given range of W, and S, a higher W, number combined with a lower 6 value leads to a maximum average aspect ratio. It can be expected that if the W, is equal to 1.0 by extrapolation, and S is in a range of 0.1-0.3, the corresponding aspect ratio may be higher, indicating

Yi

696

0.2

0' 0

I

I

I

I

0.2

0.4

0.6

0.8

d

1

normalized halved thickness of samples

Figure 17 Calculated W , number as function of the halved sample thickness.

a parameter area where the composite may be more effectively reinforced in situ by a higher aspect ratio. In this presentation, one can also find the known W , - 6 relations of the four groups of samples under different processing conditions. This three-dimensional plot exhibits a clear picture of parameters that characterize the processing-dependent fibrillation process of the TLCP phase. Therefore, it may offer guidance of practical significance for the preparation of these types of composites. Since the W&) and Ax) were known, the critical breakup time, r,, could be also calculated by:

fb(X)

Yb)

= 2t$ W.&)

(16)

here t b is the unit of time, namely t,(tb = t,). The dimensionless critical time, t g , was taken directly from the Grace criteria according to Eq. (3) as a mathematical function in the 6 range related. As shown in Fig. 10, the critical time for a drop fracture can be fitted approximately by a curve in the whole rangeof 6, both for rotational and irrotational shear. It is seen from Fig.20 that the injection flux significantly affects the critical burst time. If the samples are

*samples 2 .samples 3 +samples 4

0.01

I

viscosity ratio

Figure 18 W,

L

1

0.1

6

- 6 plot: comparison of the experimental values with the theoretical limit.


697

Figure 19 (lld) - W, - 6 relation for the processing conditions studied.

injection molded witha lower flux of 8 cm3/s, the burst time is10 times longerthan those processed with a larger flux of 80 cm3/s. Obviously, the relative smaller average degree of fibrillation in samples3 and 4, may be due to this breakup effect. In the core area the flow is more

10

E

stable, therefore, the life time of deformed droplets is longer. By plotting the rc against S, we obtained the curves reported in Fig. 21, The pattern seems similar to that by Grace, but the absolute values are significantly lower.

I

S

1

0.1

0.01

0.001

0

0.2

0.4

0.6

0.8

1

normallzed halved thlckness of samples

Figure 20 Calculated critical burst time

tc as function

of the halved sample thickness.

698

Yi

0.001

t

’

0.01

,

I

0.1 vlscosity ratlo 6

1

Figure 21 tc - S plot for the four sample groups.

The big scatter in the range of our experimental conditions and the lower values are not yet well understood.

IV. MODULUS OF DISPERSED TLCP FIBERS

A. Introduction The formation of a fibrillar structure in TLCP blends makes the mechanical properties of this kind of composites similar to those of conventional fiber reinforced thermoplastics [l 1,261. However, because the molecular orientation and fibrillationof TLCPs are generally flowinduced, the formation, distribution, and alignment of these droplets and fibers are considerably more processing-dependent. We do not know: 1. How to measure the mechanical constants of

these dispersed fibers and deformed droplets that are always embedded in a matrix. 2. If these mechanical data are materials constants or processing- and morphology-dependent. Kohli et al. [27], for instance, showed that the tensile modulus of a highly drawn PC-TLCP composite could be modeled effectively by the simple additivity rule of mixtures, while the compressionmolded composite samples with a spherical TLCP morphology had moduli according to the inverse rule. In both cases, the tensile modulus of the TLCP (EL=) itself was assumed to be a constant value determined from a tensile test of the pure TLCP samples. But whether or not the dispersed TLCP fibers and deformeddroplets have the same modulus as the bulk TLCP samples remains a question. Blizard andcoworkers [281 reported a calculated average modulus of 24.6 GPa for TLCP Vectra A950, as

microfibril dispersed in the extruded and drawn PC-TLCP composite strands, based also on the rule of mixtures. As known, this simple equation is only valid for unidirectional fiber composites. There is no parameter in the equation that is concerned with different morphological features. It is, therefore, the intention of this section to present the tensile modulus of dispersed TLCP fibers and pure TLCP bulk materials,respectively, both processed by injection molding. B.

Experimental

The polymers studied and the experimental procedure were described in the last section. The samples were made of the same PC-TLCP composite and the pure TLCP Vectra A950,respectively. To study the influence of sample forms and geometries on mechanicalconstants, plate samples(80 x 80 x 2 mm2) werealso injection molded witha film gate. Tensile sampleswere then cut from the plate in the longitudinal (//) and transverse direction, respectively. In the following text, the sampleswith their cross-section of 10 X 2 mm2 are called “thin” samples, whereas the standard tensile samples are called “thick” samples. The injection and holding pressures were 100 MPa and 50 MPa, respectively. For thick (DIN 53455 Form 3) and thin (cut from the plates) dumbbell-shaped samples, an average shear rate of 230 S - ’ was keptconstant.

c/-)

C. ResultsandDiscussion Table 4 shows that tensile modulus Et and strength U b decreased with increase of the melt temperature T,,, and shear rate y . From fractured surfaces, TLCP domains


were found by microscope to be differently deformed, fibrillated, and distributed. In allcases of the processing conditions, TLCP domains were well dispersed and deformed to droplets in the core layer, but there was onlya narrow distribution of their aspect ratio (about Ild 5 6) and less orientation. In bothtransition and skinlayers, the domains were also well dispersed, but more oriented and fibrillated in the flow direction. From this reason, we give the distribution of aspect ratio (Nd)and fiber number (N)versus fiber lengthclass in Fig. 22, only for skin and transition layers, respectively.

skin layer

699

From the difference of these morphological characteristics, a simplified three-layer model could be set up and the inhomogeneousstructure could be classified into these three layers. Figure 23 illustrates the relative thickness of the core, transition, and skin layer for the four sample groups, depending on the processing condition. Within the conditions, the extremely thick skinlayer was produced with a lower T,,, of 280°C combined with a lowery of 116 s-l. In this case, the major cross-section of these samples was filled with highlyoriented TLCP fibers. This led to the highest composite modulus and strength of 4922 MPa and 112 MPa, respectively. In-

trans. layer 0

'

Vd 30

0

tqo

40

20

40

length class Cm)

90

1-

60

0

trans. layer

skin layer 9

0

5

N 30

0

800

40

80

20

40

samples group 1

trans. layer

skin layer

skin layer

trans. kyer

60

Vd 30

0 0

40

80 0

20

40

0

40

80 0

20

40

length clam Cm) samples group 4

Figure 22 Distribution of fiber aspect ratioIld and fiber numberN versus fiber length class for skin and transition layer

of the four groupsof samples injection molded.

Yi

700

No.of the sample Trn('C) Y (Vs) sample center

31158

4 1158

sample edge (100%)

(0%)

280

.";..

,

,

.,

: r ,

320

core layer transition layer skin layer

Figure 23 Distribution of average layer thickness within a sample in dependence on melt temperatureT,,, and shear rate y by injection molding.

creasing the melt temperature and shear rate caused a decrease of degree of the fibrillation, and, therefore, the thickness of the skin and transition layer, due to the breakup time limit. This behavior is believed to be responsible for the lower mechanical values. According to the composite theory, tensile modulus of fiberreinforced composites can be calculatedby knowing the mechanical constants of the components, their volumefraction, the fiber aspect ratio, and orientation.But in the case ofin situ composites injection molded, the TLCP fibrils are developed duringthe processing andare still embeddedin the matrix. Their modulus cannot be directly measured. To overcome this problem, a calculation procedure was developedto estimate the tensile modulus of the dispersed fibers and droplets as following. The three layers can be treated as an iso-strainparallel model in the flow direction. The total cross-section Atotaland the total tensile modulus Etotal (Et)of a composite sample can be determined as follows:

where Ei stands for the moduli of the three different layers in the sample (i = 1 - 3), while Ai refers to the corresponding areas in the cross-section, respectively. On the cross-section, the thickness of each layer in the sample width and thickness direction can be assumed to be the same. The effect of the edge areas cannot be considered. According to the Cox-Darlingtonmodel [29,30], modulus Ei of each layers is expressed as below:

6

tanh x

where Ef is the modulus of dispersed fibers, namely the ELCto be determined in our case. The modulus of matrix, E,,,, is in our composites equal to 2500 MPa, 4i represents fiber volume fraction in each layer, and nr is a so-called fiber length modifyingfactor that depends on the fiber aspect ratio (Ild) and volume fraction. Where hj is equalto Nj/Ni, where Nj is the number of fibers at j class in each layer,j = 1to 6 , and Ni is the total number of fibers at layer i (i = 1,2,3), And U,,, represents Poisson's ratio of matrix, here U,,, = 0.35. Equation (18) is an implicit function ofE f , which is given in Fig. 24 by a computer simulation. It indicates that E, increases monotonously with increasingE f . The Ef can be determinedby replacing the experimental results of E, (Table 4) of the composite samples in the function. It is interesting to note that the so-obtained four moduli were found approximatelyto be a constant of 24 GPa, independent of the deformational difference of the TLCP domains. Inother words, the elastic tensile properties of the in situ formed TLCP may be changed less by varying the flow condition during the injection molding. The modulus valueseems to be a material constant of these fibers and droplets. For the calculation, the values of fld and were take from Fig. 12. These values were an average of the

701


3 6500 a

3

5500

0

4500

3500

2500 5000

1woo

25000

3SpOO

Modulus of LCP fibers E, (MPa) Figure 24 Total modulusE, of PC-TLCP composite sample as function of the fiber modulus Ef for the four sample groups.

distribution functions in Fig. 22 combined with the results in Fig. 23. The tensile modulus calculated previously isin fair agreement with that reported in the literature. Blizard et al. [28] also extruded pure TLCP Vectra A950 and its mixtures withPC witha single-screw extruder equipped with a capillary die having a diameter of 3.17 mm and a length of 9.6 mm to obtain the pure and composite strands, respectively. The extrudates were then taken up at different speeds after a quenching. Because of this stretching operation, the TLCP domains were mostly fibrillated in allcases of the composite strands and orientednearlyunidirectional in the machine direction. They reported a tensile modulus of 29 GPa for the extruded pure TLCP strand and an average modulus of 24.6 GPa for the individual TLCP microfibril by a calculationbasedon the additivity rule of mixtures. As known, the rule of mixtures is basedon the assumption of continuous fibers and therefore establishes an upper limit. Table 5 compares the tensile properties of Vectra A950 in the form of dispersed fibers and droplets in the matrix by injection molding, microfibrilextrusion by and drawing [28], injection molded pure thick sample and pure thin sample, and the pure drawn strand [28]. As exhibited, our calculated fiber modulus with its average than that of the thick and thin of 24 GPa is much higher pure TLCP samplesinjectionmolded. It can be explained that in cases of pure TLCP samples the material may only be fibrillated in a very thin skin layer owing to the excellent flow behavior in comparison with that in the blends. However, this modulus value is lower than that of the extruded and drawn pure strand. This can be

viewed as a result of the drawing. As is well known, drawing results in a high orientation and fibrillation. Fiber spun Vectra has been, for example, reported to have moduliof 41GPa at a draw ratio of 6 [31] and more than 62 GPa at draw ratios of 50 or more [32]. It appears to be natural that our calculated modulus was just between the twolimits of the pure samplesinjection molded and extruded/drawn. This constant value can be taken in turn into the composite functions to calculate the composite properties. A calculation result is illustrated in Fig. 25. For the four sample groups, the calculated layer moduli Eiare uneven in the cross-section within a composite sample group. The lowest value is stilllocated in the core layer due to the lower deformation and, therefore, the lower

Table 5 Tensile Properties of Vectra A950 as Fibers and as Pure Bulk Materials

Tensile properties TLCP fibers in injection molded thick composite sample TLCP fibrils in extruded and drawn composite strand [28] Pure TLCP thick sample Pure TLCP thin sample(//) Pure TLCP thin sample ( - L ) Extruded and drawn pure TLCP strand [28]

Modulus Strength (GPd

W")

24.0

?

24.6

-

13.5

175 130

9.5 1.7

28

53

-

702

Yi

3

1-

I

2

I

I

l

perimentally demonstrated, and a generalized function has been derived to evaluate the relationship between the blend viscosityreduction and TLCP weight fraction. Interfacial slippage of the TLCP-rich skin layer is the mechanistic explanation of the observed effect. Computer fitting and calculation of the experimental data can give property-related parameters characterizing the change manner of the blend viscosity with the TLCP weight fraction. Comparison of the theoretical curves with the experimental data gave satisfactory results. Therefore, this function has been proved to be a useful tool in dealing with the viscosity reduction behavior. Its applicability has been found to be independent of the special chemical nature of the materials. The fibrillation behavior of a dispersed TLCP domains in a thermoplastic matrix can be described by using the W, - S relation initially given by Taylor and Cox. The flow-induced deformation istime-dependent. If the flow time exceeds the critical time, the resulting fibers break up. The life time of a TLCP fiber can be evaluated by Grace criteria. Based on morphological and rheological examinations, a calculation procedure was developed to predict the fibrillation behavior of a TLCP phase in a thermoplastic matrix during injection molding.In particular, the W, - S - l/d relation (Weber number of the melt, respectively for W, - S - Id, viscosity ratio of TLCP to matrix aspect ratio of the fibrils) may give complete information for processing optimization in order to achieve highly oriented TLCP fibrous structure. The variation in the tensile modulus of a dispersed TLCP phase depends primarily onthe degree of orientation and fibrillation,or in other words, on the fiber-aspect ratio. Injection molding or extrusion offer samples that give only an integrated value. It can not be taken as the modulus of the in situ formed TLCP domains directly. In the case of injection molded blend samples, the modulus of dispersed TLCP domains can be estimated bya calculation procedure described in this chapter. For Vectra A950 tested, as it dispersed and deformed in the PC matrix, the average tensile modulus was found to be approximately constant, independent of the processing conditions. For most composites containing a deformed inclusion, the proposed calculation procedure may offer a useful wayto estimate the modulus of the inclusion that is usually unknowndue to testing problems.

1 E"_

'

" " " " " " " " " "

t""" I

I

I

0

I 100%

Halved sample thickness

Figure 25 Calculated layer modulus Eias function of the halved sample thickness for the four sample groups anda comparison with the total modulus Et.

aspect ratio of the TLCP droplets. In the transition layer, the modulus obtains its maximum. The experimentally measured total modulus E, is approximatelythe average of Ei. Thus, the calculation procedure has been proven as valid for the in situ composites to estimate the modulus values.

V. CONCLUSIONS A distinct viscosity reductionoccurs as a small quantity of TLCP is added to a thermoplastic matrix. This negative deviation effect from the log-additivity rule was ex-

ACKNOWLEDGMENTS

This study was supported by the National Natural Science Foundation of China (NNSFC, Grants No. 58903427 and 59183016),the National Advanced Materials Committeeof China(NAMCC, Grant No. 86371527), and the Deutsche Akademische .Austauschdienst (DAAD, Germany) Foundation.

Thermotropic

703

layer cross-section in three layers model (i = 1-3) total cross-sectionof a tensile sample width of a dumbbell-shaped sample layer thickness, average fibril diameter, short axis of a deformed droplet tensile modulus modulus of dispersed TLCP fibers (f) and of matrix (m) layer modulus in three layers model (i = 1-3) modulus of a composite sample experimentally obtained thickness of a dumbbell sample linear factor long axis of a deformed droplet fiber aspect ratio interlayer number within a capillary flow, flow index of viscosity curve fibril number withina layer fiber length modifying factor number of fibers at layer i(i = 1 - 3) fiber number a t j class in each layer ( j = 1 - 6) volume flux distance of a TLCP-rich interlayer to the capillary axis critical radius capillary radius breakup time dimensionless breakup time melt temperature interface slide velocity Weber number thickness variable interface slip factor proportional factor shear rate viscosity ratio of ql/qO exponent blend viscosity matrix viscosity or qc viscosity of continuous phase (qc = qo) TLCP viscosity or 770 viscosity of dispersed phase (70 = q l ) critical stress ratio Poisson’s ratio of matrix density interfacial tension coefficient tensile strength critical interface shear stress shear stress shear stress on the capillary wall

4

= weight fraction of TLCP

4i

= fiber volume fraction in each layer

(three layers model)

REFERENCES 1. A. Siegmann, A. Dagan, and S. Kenig, Polym.. 26: 1325 (1985). 2. D. Dutta, H. Fruitwala, A. Kohli,and R. A. Weiss, Polym. Eng. Sci., 30: l005 (1990). 3. M. R. Nobile, E. Amendola, L. Nicolais, D. Acierno, and C. Carfama, Polym. Eng. Sci., 29: 244 (1989). 4. L. A. Utracki, M.M. Dumoulin, P. Toma, Polym. Eng. Sci., 26: 34 (1986). 5. N. Chapleau, P. J. Carreau, C. Peleteiro, P. A. Lavoie, and T. M. Malik Polym. Eng. Sci., 32: 1876 (1992). 6. A. I. Isayev and M. J. Modic, Polym. Compos., 8: 158 (1987). 7. X.-S. Yi, L. Wei, and H. Wang, Cailiao Kexue Jingzhan, (Chin. J. Mat. Resear.), 6:256 (1992), (Chinese). 8. X.-S. Yi, Preparation and Processing of Plastics, Zhejiang University Press, Hangzhou(1995) (Chinese). 9. V.G. Kulichiklin, 0. V. Vasil’eva, I. A. Litinov, E. M. Antopov, I. L. PasamyanandN. A. Plate, J.Appl. Polym. Sci.. 42: 363 (1991). 10. F. D. Rumscheidt and S. G. Mason, J. Colloid Sci., 16: 238 (1961). 11. X.-S. Yi,’ G. Zhao, and F. Shi, Polym. Intern 39(1): 1 1 (1996). 12. X.-S. Yi andG. Zhao, J. Appl. Polym. Sci., 61: 1655 (1996). 13. X.-S. Yi, andL. Shen, Po1ym.-Plast. Technol. Eng. 36(1): 153 (1997). 14. J.-C. Huang, H.-F. Shen, and Y.-T. Chu, Adv. Polym. Techn., 13: 49 (1994). 15. L. A. Utraki and P. Sammut, Polym. Eng. Sci.,30: 1027 (1990). 16. G. I. Taylor, Proc. Roy. Soc. (London), Ser. A146: 501 (1934); and Ser. AI38: 41 (1932). 17. C. D. Han, Multiphase Flow in Polymer Processing,Academic Press, New York (1981). 18. H. P. Grace, Chem. Eng. Commun., 14: 225 (1982). 19. J. J. Elmendorp, Polym. Eng. Sci., 26: 418 (1986). 20. M. V. Tsebrenko, G. P. Danilova, and A. YA. Malkin, J. NonrNewton. Fluid Mech., 31: 1 (1989). 21. B. J. Bentley andL. G. Leal, J. Fluid Mech., 241 (1967). 22. N. Chapleau, P. J. Carreau, C. Peleteiro, P. A. Lavoie, and T. M. Malik, Polym. Eng. Sci., 32: 1876 (1992). 23. Th. Brinkmann, PhD. Thesis, RWTH Aachen (1992). 24. J. Schmitz, E. Bornschlegel, G. Dupp, and G. Erhard, Plastverurbeiter 39: 50 (1988). 25. R. G. Cox, J. Fluid Mech., 37(3):601 (1969). 26. G. Kiss, Polym. Eng. Sci., 28: 1248 (1988). 27. A. Kohli, N. Chung, and R. A. Weiss, Polym. Eng. Sci., 2 9 573 (1989). 28. K. G. Blizard, C. Federici, 0. Fererico, and L. L. Chapoy, Polym. Eng. Sci., 30: 1442 (1990). 29. X.-S. Yi and L. Shen, Gaofenzi Xuebao (Acta. Polym. Sinica), 5: 621 (1994) (Chinese). 30. H. L. Cox, Brit. J . Appl. Phys., 3: 72 (1952). 31. S . Kenig, Polym. Eng. Sci., 2 9 1136 (1989). 32. G. W.Calundann andM. Jaf€e, Proc. Robert WelchCon$ XXVZ, Synthetic Polymers, 247 (1982).


Energetic Composites M. B. Khan* CHEMTEC & Prime Glass, Jhelum, Pakistan

1. INTRODUCTION Energetic composites, commonly referred to ascomposite propellants, are ignitable assemblies containing energetic particles (inorganic oxidizer and metal) that are embedded ina viscoelastic matrix. The primary function of these composites is to supply the energy needed for the propulsion of space vehicles and tactical missiles. The composite is normally contained in a casing, lined with an ablative thermal insulationto provide protection from the ultra-high temperatures (3OOO-3600°K) encountered in the combustion chamber. In the missile motor (Fig. l), the sensitive heat available from the combustion reaction isconverted into the kinetic energy of the gases at the nozzle exit plane, providing the thrust needed for vehicle propulsion. From the ingredients to feed preparation to final fabrication in the casing, the solid fuel composite passes through several manufacturing stages, during which itis strictly scrutinized to achieve the requisite mechanical and ballisticproperties in the end. Unlike other strategic technologies, energetic composites is a well-published discipline. The analytical and experimental efforts that have beenundertaken to formulate, fabricate, characterize, and preserve these systems for long-term storage have all been welldescribed. While abundant literature is available onthe various aspects of the subject [l-lo], a comprehensive treatment from the materials processing standpoint seems to be lacking. Processing constitutes an important sector in the production of energetic composites. It is during this phase that new material interfaces are formed, and the viscous prepolymer slurry is converted to a viscoelastic composite. This chapter discusses the technical features of the production sequence, within the general domain of vis-

Casing Ablative insulation

\

-\(ExpanJon nozzle Energetic Composite Figure 1 Energetic composite assembly showing various

components.

cous reactive processing, with special emphasis on multiphase rheology,the chemorheology andthe associated morphology of the reactive polymer system, interfacial phenomena, and the equipment configuration. The discussion largely relies on the elucidation of the governing principles, withrecourse to experimental data wherever necessary. The main discussion is preceded by a brief overview on some of the key performance parameters that influence the choice of the basic constituents. The chapter concludes with a technical description of the principal processing equipment used in the production of energetic composites. 11.

OVERVIEW

A. Ingredients An energetic composite is basically a fuel oxidizer assembly containing several important additives to perform specific functions. The fabricated system derives

* Former consultant, Pakistan Air Force. 705

706

Khan

its propulsion force from the combustion reaction between the oxidizer and the fuel components. The oxidizer together with the various ingredients are incorporated in a hydrocarbon fuel, a processable liquid prepolymer that upon curing transforms into an elastomeric matrix.The composition of the system is dictated by the often conflicting ballistic and processing constraints. To meet only the minimum performance requirements, the solid content (oxidizer plus metal) the in composite should be greater than 80%. This high level of solids loading limits the flexibility of the chemist in adjusting the content and rheology of the composition. Thus, matching performance and processability is often the real difficulty. Aside fromthe oxidizer component, metal particles are added to augment the heat of combustion; ballistic catalysts may be used to enhance the burning rate; crosslinking and interfacial agents may serve to improve the mechanical properties; andpolymerization catalysts may help to control the formation of the crosslinking network at a predetermined rate. Plasticizers and wetting agents are almost always used to aid inthe processing, as they tend to reduce the viscosity of the mix to acceptable levels for the subsequent pouring andcasting operations. Finally, the addition of an antioxidant improves shelf-lifeby arresting bond cleavage and/or crosslinking reactions at sites of unsaturation in the otherwise crosslinked polymer.

B. MaterialandPerformanceParameters In formulatingenergetic composites, the primary aim is to pack maximum thermal energy per unit volume. A

principal performance characteristic of the fuel is the specific impulse, defined as thethrust delivered per unit mass rate of fuel burnt. From the ballistic standpoint, it is desirable to maximize the specific impulse of the propulsion unit, simply expressed as:

zsp = vJg

=

f(m

(1)

The solid fuel compositionis, therefore, concerned with the pursuit of higher Zsp involving both flame temperature elevation and molecular weight reduction. Concomitantly, it is desirable to have high heat of combustion, which also dictates the combustion chamber temperature. The actual composition is, therefore, set on the basis of the energy quantuminherent in the various constituents. The fuel combustion enthalpy is primarily determined bythe polymer binder and the metal heats of combustion. The polymer, which binds the particulates together, is an important constituent of the reactive composite: it contributes to the overall enthalpy to the tune of 30-40 x IO' kJ/kg, apart from conferring vital mechanical properties, such as tensile strength and elastic modulus. A1 is the preferred choice as the metallic fuel (30 MJ/kg) because of its ease of incorporation and much higher specific gravity (2.70) compared withother more thermally active metals such as Be (63 MJ/kg and 1.81). A higherspecificgravityis synonymous with higher massratios, which augmentsthe vehicle burnout velocity. Modern propellants use polybutadiene type binders whose main chain terminates with either hydroxyl

(HTPB: HO-[CH2"CH=CH"CH2).dCH+2H).z"(CH&H"CH2).&&H JH=CH2

or carboxyl (CTPB: HOOCR-(CHflH=CH-CH2)nend groups. For obvious reasons, polymers containing

a high hydrogen content and higher heats of formation (those absorbing less heat on combustion) yield more energeticcompositions. A comparison is shown in Table 1 between two commonly used prepolymers, namely HTPB and PPG (polypropylene glycol) [9]. The higher heat of formation andless oxygen in the HTPB molecule implies a higher heat output with greater oxidizer loadingcapacity. However, more oxygenin the

PPG would allow more metalat the expense of the oxidizer to further energize the composition. In addition, HTPB is frequently preferred over PPG because of its superior mechanicalproperties, better aging characteristics, and lower glass-transition temperature (Tg). The latter is especiallydesirable because at low temperature, the higher strain rates produced by motor ignition decrease the elongation of the composite rather markedly. C.

Table 1 ComparisonofPrepolymer

Heats of Formation

Content Oxygen and

A H ~ ( I LXT 103) ~~ HTPB PPG

20.92 - 3740

02%

5.0 26.0

RCOOH)

Composition

Several techniques are availablein the literature for evaluation of the flame temperature, exit temperature, equilibrium composition of combustion products, and performance parameters of energetic composites [l 1-13]. The optimum combination of the composite indetermined gredients is by thermodynamic means, so as to arrive at a composition having maximum performance

707

Energetic Composites

capabilities. The optimization procedure is, however, no longer dependent on trial and error methods. Thermochemical computer codes are now available [14,15] in which the motor thrust is expressed as a function of the reactants and optimized subject to linear or nonlinear constraints. Table 2 lists the various ingredients along with the possible rangeof composition for a typical energetic composite. D.

PerformanceEnhancement

The major emphasis the in development of new formulations has been on composites where an energetic solid is dispersed in an energetic polymer matrix. The polymer system in contention is the glycidyl azide polymer (GAP). Its curing properties are not ideal, but it was really designed to demonstrate that the energetic polymer approach is feasible[ 161. Other advanced high-energy systems use nitramine (RDX or HMX) crystals in the so-called crosslinked double-base (XLDB) propellants, where a conventional prepolymer (PPG, HTPB, CTPB, PBAN, etc.) is mixed with double-base ingredients (nitrocellulose and nitroglycerine) together with a relatively largeamount of RDXor HMX andcrosslinked to provide a highly energized, HC1-free system, which also obviates a major ecologicaldrawback of AP (ammonium percho1orate)-based composites [17]. The inclusion of energetic solidsinhigh-energy polymers also increases their vulnerability to detonation. The search, therefore, continues for new energetic solids that exceed the energy and density of RDX, but with less thermal and shock sensitivity. Polycyclic and bridged nitraminesare families of new energetic materials with much promise.Plasticizers are also key ingredients because of their influence on energy, mechanical properties, and processing characteristics. Plasticizers containing nitric ester groups combine the ease of processing with higher energy content, but pose compati-

Table 2 Composition Range of a Typical Energetic

bility and sensitivity problems. A promising approach is the use of molecular modeling software to elucidate polymer-plasticizer interactions and evaluate the potential of newcompounds. New and emerging materials exploit the nitramino and azado moieties for superior energy and reducedsensitivity relative to nitrato moieties [N.

Finally, there is active interest in developing catalyst systems, bothballisticand polymerization, that would promote combustion stability at high pressures (especially in metal-freesystems for smokeless applications) and allow processinglattitude for relatively large motors. The ferric-based systems currently being used fall short of these performance measures. Compounds that form complex structures with the metal chelate to reduce its activity to acceptable levels seem to be most promising. Interestingly, the use of an antibiotic has been cited in this context [19]. 111.

CONCENTRATEDPARTICLE PROCESSING

A.

MultiphaseRheology

Rheology is the science of flow anddeformation of matter. It is of direct importance to theproduction of energetic composites; processability heavilydepends on the flow properties of the composite slurries. The process fluid in these systems is of a multiphase character, with a high percentage of solids suspended in a viscous and reactive liquid. The degree of particle loading, their size distribution, interfacial effects, and polymerization may substantially influence the rheological behavior of the system. However, the primary parameter that determines the bulk viscosity is particle concentration. This section will describe the rheology of suspensions and highlight the nature of particle assemblies that may be accommodated in the liquid, according to a particular size distribution. The influence of polymerization kinetics on the rheology will be discussed in a subsequent section.

Composite

l . Estimation of Suspension Viscosity

Ingredient Oxidizer Metal fuel Prepolymer Polymerizer Crosslinker Chain extender Plasticizer Bonding agent Antioxidant Ballistic catalyst Curing catalyst

Perchlorate; Nitramine Al; Be; BeHz Polyglycol; PB NCO; Epoxide Triol Diol Alcohol derivative Polyalkylamine PhenoVamine derivative Metal oxide Metallic derivative

60-88 20-40 0-2 1 8-12 1-2 0-3

0-3

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

Trace

An estimation of the multiphase viscosity is a preliminary necessity for convenient particle processing. For particle-dopedliquids the classical Einstein equation [20] relates the relative viscosity to the concentration of the solid phase: 7, = 1

+ 2.54

(2)

Equation (2)is valid only for very dilute suspensions of nondeformable, smooth, uniform spheres. It assumes a Newtonianliquid phase and neglects interaction between particles, a plausible condition when the volume of the solid phase issmallcomparedwith the liquid phase.

708

Khan

For relatively higher volume fractions (250°C and are, therefore, not usable as a thermoplast matrix for natural fibers [60]. Other than the processes mentioned here, natural fibers are used as construction units, by applying hybrid nonwovens, i.e., natural fiber staple fiber fleece [61].

' I . NaturalFiberMat-Reinforced Thermoplastics (NMT) With the largest turnover, GMT is certainly the most important semi-product inthe group of reinforced thermoplastics. A specialproduction process for natural fiber-reinforced PP semi-products (NMT) has been de-

veloped by BASF AG[60]. For this process, natural fibers must be availablein form of fiber mats. Mats are produced by stitching together layers of fibers that have been crumbled. Continued production of this semi-product is done by melt-coating the fiber mats in a double coil coating press, which is furnished with a heat- and cool-press zone. In such a coil coating press, the fiber mats are brought together with the polypropylenic melt between circulating steel bands. Wetting of the mats with the thermoplast melt takes place in the hot-press zone. Then the laminate is cooledunder pressure in the cooling-press zone.

60

F 5

40

I

5 L

c

U1

a,

420 c

F

0 0

20

40 80

60

100

120

Fiber content [vol%]

Figure 18 Tensile strength of B.M.C. molded plastics as dependent on their fiber content. %fibers = g fibed100 g UP

resin.

Natural Fiber Reinforced Plastics

805

Table 12 Tensile Strength and Young’s Modulus

nonwovens on the bottom and one on top, and in between the melt-film. In reference to density, the values for tensile strength of flax fibers PP composites exceed the characteristic values of sisalfiber-reinforced PP composites (Fig. 19).

of Sisal, Flax, and Glass Fiber MTs with a Fiber Content of 40% (weight) 1601

Materials PP-Wood flour 3600 PP-Sisal nonwoven PP-Sisal nonwoven with surface treatment PP-Flax nonwoven 6700with PP-Flax nonwoven surface treatment PP-Glass nonwoven

Tensile strength (MW

Young’s modulus (MPa)

19

2500

55

4800

47 67

5100

100

m

38

D. Influence of FiberContent on the Mechanical Properties of Natural Fiber-Reinforced Plastics

Sisal, flax, and glassfiber MTs can be classified by their mechanical properties, tensile strength, and Young’s modulus (Table 12). 2. “Express”Processing Extrusion press processing (express processing) was developed for the production of flax fiber-reinforcedPP at the research center of Daimler Benz (Ulm, Germany) [62]. In this processing, natural fiber nonwovens and thermoplastic melt-films are alternatively deposited in a tempered molding tool and molded afterwards. The thermoplasticmelt-films are laidonby a mobileextruder. If this process is optimally adapted to the element, a single passage by the extruder suffices. The structural order consists of three layers: two layers of

120 I

As is known of glass fiber-reinforced plastics, the mechanical and physicalproperties of composites, next to the fiber properties, and the quality of the fiber matrix interface, as well as the textile form of the reinforcement primarily depend on the volume content of fibers in the composite. jute fiberTests by Roe et al. [63] with unidirectional reinforced UP resins show a linear relationship (analogous to the linear mixingrule) between the volume content of fiber and Young’s modulus and tensile strength of the composite over a range of fiber content of 0-60%. Similar results are attained for the work of fracture and for the interlaminate shear strength (Fig. 20). Chawla et al. [M] found similar results for the flexural properties of jute fiber-UP composites.

E. Hybrid Composites Made of Natural Fibers and Glass or Carbon Fibers

Generally, the mechanical and physical properties of natural fiber-reinforcedplastics only conditionallyreach the characteristic values of glass fiber-reinforced systems. By using hybridcomposites made ofnatural fibers and carbon fibers or natural fibers and glass fibers, the

1

I

l

EXPRESS glass

MPRESS

l00

S

E.

MPRESS flax

80

0

5

U)

g -ag .c

/

4’

60 EXPRESS

(I)

/ sisal 40

F

\

20

I

I

PP-Matrix

01

099

1 9 1

1,3

II

1 3

Density of the construction unit [g/cm3]

Figure 19 Tensile strengthof different types of fiber reinforced composites, produced by the express-processing 1621.

Bledzki and Gassan

806

~

0

0,1

0.2

0.3

0.4

0.5

0.6

Volume taction

0.1

0.2

03

0.4

Volume fraction

0.5

0,s

0.7

Volum fraction

Volumo fraction

Figure 20 Influence of fiber contentby volume on tensile strength,Youngs modulus, work of fracture, and interlaminate shear strength of one-dimensional jute fiber-reinforced UP resins [63].

properties of natural fiber-reinforcedcomposites can be improved further [65-691, as Fig. 21 exemplifies for compression strength. To this, examinations were made by Mohan et al. [67,69], Philip [70], and Clark et al. [68] on jute-glass hybrid composites; Pavithran et al. [65] on coir-glass hybridcomposites; Chand et al. [66,71,91] on sunhemp-carbon hybrid composites. Additionally, the dependence of the mechanical properties on humidity is clearly reduced because of the moisture repellency of glass fibers [69]. The moisture absorption of composites is clearly smaller when natural fibers are replaced by glass fibers [65].

F. Biologically DegradableComposite Materials

For ecological reasons, more activities in the area of biologically degradable composite materials, i.e., natu-

ralfiber-reinforcedbiologically degradable polymers, were recognized. Appliedas a matrix are the most readily available polymersas Biopol, PHB-HV coploymer, Bioceta, Mater Bi, or Sconacell A [72]. Tests with different flax fiber-reinforced biologically degradable polymers (Table13) by Hanselka et al. [721 show that tensile strength and Young’s modulusof these composites are clearly influenced bythe particular matrix andthe adhesion between fiber and matrix. Tests by Gatenholm et al. [8,10] on PHB-HV copolymers containing cellulose fibers (for example, the tradenamed Biopol) show that the mechanical properties of these systems are determined by the fiber and the fiber matrix interface on the one hand, and on the other hand by the composition of the matrix, that is, of HV proportion in the matrix. At an increased proportion of HV, the stiffness of the composite is reduced upto 30%, whereas elongation at break increases until about 60%.

P@


807

0

0.2

JRC

0.4

0.6

Volume fraction of GRP

l.Q GRP

Figure 21 Compression strength as dependent on the content of GRP in jute fiber-reinforced hybrid-composites [67].

Table 13 Mechanical Properties of Flax Fiber

Reinforced Biologically Degradable Polymers (BDG)[72] Tensile strength

Fibers-BDP

Young's modulus W")

65.7 106.9 124.3

Flax-Bioceta Flax-Sconacell A Flax-Mater Bi

1,400 8,180 10,580

Fiber content = 50% by vol., linen weave, cable direction= testing direction.

But the processing parameters are being greatly influenced by the proportion of HV. Tests made by Avella et al. [73] on straw fiber-reinforced PHB generally lead to expect good mechanical properties of such composites. The mechanical properties of extruded flax fiberreinforced thermoplastic starch (structured with water or glycerin), clearly show increased values, especially for tensile strength and Young's modulusbecause of the addition of green-flax fiber rovings. Experiences mentioned in the literature on this subject suggest that from the point of view of the mechanical properties such biocomposites are suitable construction materials. At pres-

"

0

4

8

Percent fiber content [by mass]

12

Percent f l b r content by mass]

Figure 22 Influence of fiber content on flexural strength and fracture toughnessof (0)softwood-cement composites and

(0)hardwood-cement composites (air-cured)[78].

r.h.

808

Bledzki and Gassan

Table 14 Influence of Humidity on Flexural Strength

and Fracture Toughness of Cellulose Fiber Reinforced Cements [78] Fracture

Chemicals/time

toughness Flexural Fiber (kJ/m2) (MPa) strength content (wt%) 2 4 6 8 10 12

20.9 20.3 20.1 20.6

8.6 10.5 10.4 8.4 9.6 9.3

0.25 0.51 1.06 1.37 1.46 1.68

Table 15 Influence of Chemicals on the Change of Bending Strength of Sisal-PP Composites 1791

0.33 1.oo

1.61 1.49 1.83 1.79

r.h. = test samples were conditioned before testing at 50 k 5% r.h. and 22 f 2°C; wet = test samples were soaked in water for 48 h, dried with a cloth and then tested while wet.

ent, limitations must be seen where excessive environmentalconditions exist. Major prospects for these material systems are, therefore, lining elements with support function in the automobile, rail car, and furniture industries [72]. G. AdditionalApplications

When usedas substitutes for asbestos fibers, plant fibers and "manmade" cellulosefibers show comparable characteristic values ina cement matrix, but at lower costs. As with plastic composites, these values are essentially dependent on the properties of the fiber and the adhesion between fiber and matrix. Distinctly higher values for strength and stiffnessof the composites can be achieved by a chemical modification of the fiber surface (acrylic and polystyrene treatment [74]), usually producedby the Hatschek-process [75-771. Tests by Coutts et al. [76] and Coutts [77,78] onwood fiber cement (soft-, and hardwood fibers) show that already at a fiber content of 8-10 wt%, a maximum ofstrengthening isachieved (Fig. 22). Cellulose cement composites show a similar sensibility to humidity (Table 14), as do plastic composites, that is, they show a decrease of mechanical properties. Yet, formerly reached values can again be achieved by a drying process on the composite [75]. V. ENVIRONMENTAL EFFECTS Whichever applicationof natural fiber or natural fiberreinforced plasticswill be useddepends on the different environmental conditions, which are likely to add to the aging and degrading effects. On the other hand, such effects are often desirable, as is the case withcom-

BendingstrengthBendingmodulus W") (GP4

Reference samples NaOH/5O h NaOH/SOO h HCV50 h HCV5OO h

30 24 18 20 15

1.5 1.1 1.05 1.35 1.4

postable materials.Natural fibers are subject to degradation in acids and in alkaline solutions, as well as under UV rays. These effects, however, can be minimized by using suitable modifications. Unmodifiedcellulose fiber are normallydegradedbyenzymes after about 6-12 months; this can bealtered through suitable treatments, so that no significant changes of mechanical properties can be noticedfor 2 years. Within a period of 2.5 years, dry-stored fibers show only little changes in their mechanical properties. This is especially true with regard to strength and elongationat break. In this respect, sisal fibers are comparably more stable than Henequen and Abaka. Lower temperatures, such as -70°C, clearly result in lesser strength, but this effect can be minimized by previous drying. Higher temperatures, such as 100-13OoC, lead, in the case of cotton, to a noteable degradation after 80 days. Their strength is thereby reduced to68%, that is 10% of the original value. Depending ontemperature applied, these values are reduced to 41% and 12%, for flax fibers, and to 26% and 6%, for ramie fibers. In composites, moisture content results in a decrease of mechanical properties. This effect is greater with seawater than with freshwater. With freshwater, it is more likely that bacteria and fungi will appear. Against such influences, ramie, jute, and kopakfibers, are more resistant than other plant fibers [l]. Said lowerresistants of natural fibers, against environmental factors, decisively effects the mechanical properties of the composites (Table 15) [79]. VI. CONCLUSIONS The mechanical and physical properties of natural fibers vary considerably, as it is withall natural products. These properties are determined by the chemical and structural composition, which depend on the fiber type and growthcircumstances. With thiscellulose, the main component of allnatural fibers varies fromfiber to fiber. The moisture sensibilityis remarkable, certain natural fibers can easily be influenced byenvironmental effects. Generally speaking,the mechanical properties are lowered with rising moisturecontent.


Natural fibers compete with technical fibers, such as glass fibers or carbon fibers, as reinforcements for plastics. The advantages of technical fibers are their good mechanicalproperties, which vary onlya little, but their recycling is difficult. The mechanical properties of composites are mainly influenced by the adhesion between matrix and fibers of the composite. As it is known from glassfibers, the adhesion properties could be changed bypretreatments of fibers. So special process, chemicalandphysical modification methods were developed. Moisture repellency, resistance to environmental effects, and, not at least, the mechanical properties are improved by these treatments. Various applications for natural fibers as reinforcement in plastics are encouraged. Several of these natural fiber composites reach the mechanical properties of glass fiber composites, they are already being used in the automobile and furniture industries. Up to now, the most important natural fibers were jute, flax, and coir. Yet, the development of processing and modification methodsis not finished.Further improvements need to be expected so that it might be possibleto substitute technical fibers in composites even more widely. Natural fibers are reusing raw materials and they are recyclable. When recognizing the need for recycling and preserving natural resources, such a substitution is very important.

REFERENCES 1. A.K. Bledzki, J. Izbicka and J. Gassan, KunststoffeUmwelt-Recycling, Stettin (Poland)September 27-29 (1995). 2. W.Wittig, Kunststoffe im Automobilbau, VDI-Verlag, Diisseldorf (1994). 3. P. K. Pal, Plastics and Rubber Processing and Applications, 4:215, (1984). 4. A. G. Winfield, Plastics and Rubber International, 4(1): 23 (1979). 5. N. N., ECJWZD 34, p.15. 6. K. P. Mieck and T. Repmann, Kunststoffe, 85(3): 366 (1995). 7. N. N., Ingenieur-Werkstoffe4(9): 18 (1992). 8. R. Kohler and M. Wedler, Landinfo, 3: 33 (1995). 9. R. W. Kessler, U. Becker, R. Kohler and B. Goth, Biomass and Bioenergy, (in press). 10. P. Gatenholm,J. KubiitandA. Mathiasson, J . Appl. Polym. Sci., 45: 1667 (1992). 11. P. Gatenholm and A. Mathiasson, Polymeric Materials Science and Engineering, 6 7 361 (1992). 12. J. Gassan andA. K. Bledzki, Die Angewandte Makromolekulare Chemie (in press). 13. K.-P. Mieck, A. Nechwatal and C . Knobelsdorf, Melliand Textilberichte, 11: 892 (1994). 14. P. S. Mukhejee and K. G. Satyanarayana, J . Mat. Sci., 21: 51 (1986). 15. J. Gassan andA. K.Bledzki, 6.Znternationales Techtexil Symposium 1994, Frankfurt July 15-17 (1994).

809 16. G. E. Kritschewsky, Chemische Technology yon Textilmaterialien, Legprombitisdat Moskau(1985). 17. H. P. Fink, J. Ganster andJ. Fraatz, Akzo-Nobel Viskose Chemistry Seminar a ‘Challenges in cellulosic man-made fibers”, Stockholm, May 30-June 3 (1994). 18. J. Gassan and A. K. Bledzki, 7 . Internationales Techtexil Symposium 1995, Frankfurt, June 20-22 (1995). 19. J. W. S. Hearle and J. T. Sparrow, J . Appl. Polym. Sci., 24: 1857 (1979). 20. P. S. Mukhejee and K. G. Satyanarayana, J . Mat. Sci., 21: 51 (1986). 21. J. Gassan and A.K. Bledzki, 9th International Conference on Mechanics of Composite Materials, Riga (Latvia), October 17-20 (1995). 22. S. H. Zeronian, H. Kawabataand K.W.Alger, Text. Res. Inst., 60(3): 179 (1990). 23. M.A. Semsarzadeh, Polym. Comp., 7(2): 23(1986). 24. M. A. Semsarzadeh, A.R. Lotfaliand H. Mirzadeh, Polym. Comp., 5(2): 141 (1984). 25. P. K. Ray, A. C . Chakravarty and S. B. Bandyopadhyay, J . Appl. Polym. Sci., 20: 1765 (1976). 26. A. N. Shan and S. C. Lakkard, Fibre Sci. Techn., I S : 41 (1981). 27. B. Wulfhorst,G. Tetzlaffand R. Kaldenhoff, Techn. Text., 35(3):S . 10 (1992). 28. M.N.Belgacem, P. Bataille and S. Sapieha, J . Appl. Polym. Sci., 53:379 (1994). 29. I. Sakata, M. Morita, N. Tsuruta andK. Morita, J . Appl. Polym. Sci., 49: 1251 (1993). 30. Q.Wang, S. Kaliaguine and A. Ait-Kadi, J . Appl. Polym. Sci., 48: 121 (1993). 31. S. Goa and Y. Zeng, J . Appl. Polym. Sci.,4 7 2065 (1993). 32. S. Dong, S. Sapieha and H.P. Schreiber, Polym. Eng. Sci., 32: 1734 (1992). 33. T. P. Nevell andS. H. Zeronian, Cellulose Chemistryand its Applications, John Wiley & Sons, New York (1985). 34. S. C. 0.Ugbolue, Text. Inst., 20(4): 1 (1990). 35. J. I. Kroschwitz, Polymers:Fibers and Textiles, John Wiley & Sons, New York (1990). 36. V. V. Safonov, Treatment of Textile Materials, Legprombitizdat, Moscow (1991). 37. S. H. Zeronian, J . Appl. Polym. Sci., 4 7 445 (1991). 38. B. S. Westerlind and J. C . Berg, J . Appl. Polym. Sci., 36: 523 (1988). 39. M. J. Schick, Sudace Characteristics of Fibers and Textiles. Part If, Marcel Dekker, Inc., New York (1977). 40. J. M. Felix and P. Gatenholm, J . Appl. Polym. Sci., 42: 609 (1991). 41. D. Maldas, B. V. Koktaand C. Daneaulf, J . Appl. Polym. Sci., 37: 751 (1989). 42. K. L. Mittal, Silanes and Other Coupling Agents, VSP BV, Netherlands (1992). 43. R. G. Raj, B. V. Kokta, F. Dembele andB. Sanschagrain, J . Appl. Polym. Sci., 38: 1987 (1989). 44. W. Gepner, ChemiefasernlTextilind. (1nd.-Text.), 391 91(7/8):S.185 (1989). 45. A. C . Khazanchi, M. Saxena and T. C . Rao, Text. Comp. Build. Constr. 69 (1990). 46. P. Gatenholm, H. Bertilsson andA. Mathiasson, J . Appl. Polym. Sci., 49: 197 (1993). 47. P. Ghosh, S. Biswas andC . Datta, J . Mater. Sci., 24: 205 (1989). 48. L. Hua, P. Flodin and T. Ronnhult, Polym. Comp., 8(3): 203 ( 1987). 49. L. Hua, P. Zadorecki and P. Flodin, Polym. Comp., 8(3): 199 (1987).

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50. D. Maldas, B. V. Kokta andC. Daneault, J . Vinyl Techn., 11(2):90 (1989). 51. P. Zadorecki and T. Ronnhult, J . Polym. Sci.: Part A: Polym. Chem., 24: 737 (1986). 52. P. Zadorecki and P. Flodin, J . Appl. Polym. Sci., 31: 1699 (1986). 53. M. H.Schneider and K. I. Brebner, WoodSci. & Techol., 19: 67 (1985). 54. D. S. Varma, M. Varma and I. K. Varma, J . Rein5 Plast. Comp., 4(10):419 (1985). 55. E. T. N. Bisanda and M. P. Ansell, Comp. Sci. Techn. 165 (1991). 56. R. G. Raj and B.V. Kokta, Polymer Engineering and Science, 31(18): 1358 (1991). 57. K.-P. Mieck, A. Nechwatal and C. Knobelsdorf, Die Angewandte Makromolekulare Chemie225: 37 (1995). 58. R. Kohler and M. Wedler, 6.InternationalesTechtexil Symposium 1994, Frankfurt, June 17-20 (1995). 59. 0. Owolabi, T. Czvikovszky and I. Kovhcs, J . Appl. Polym. Sci., 30: 1827 (1985). 60. H. Baumgartl and A. Schlarb, 2.Symposium “Nachwach-

send Rohstoffe-Perspektiven fur die Chemie,” Frankfurt, May 5-6 (1993). 61. R. Liitzkendorf, K. Mieck andTh. Reupmann, 7.InternationalesTechtexilSymposium 1995, Frankfurt, June

20-22 (1995). 62. Th. Schloper and Th. Folster, Kunststoffe, 85(3): 319 ( 1995). 63. P. J. Roe and M.P.Ansell, J . Mat. Sci., 20: 4015 (1985). 64.

K. K. Chawla and A. C. Bastos, 3 International Conference on Mechanical Behaviour of Materials, Cambridge (England), Aug. (1979). 65. C. Pavithran, P. S. Mukhejee and M. Brahmakumar, J . Reinf. Plast. Comp., 10: 91 (1991). 66. N. Chand and P. K. Rohatgi, J . Mat. Sci. Letters, 5: 1181 (1986).

67. R. Mohan, M. K. Shridhar and R. M. Rao, J . Mat. Sci. Letters, 2: 99 (1983). 68. R.A. Carkand M. P. Ansell, J . Mat. Sci., 21: 269 (1986). 69. R. Mohan and J. Kishore, J . Reinf. Plast. Comp., 4: 186 (1985). 70. A.R. Philip, Engineering Materials and Design, 8: 475 (1965). 71. N . Chand and P. K. Rohatgi, Polymer Communication, 28(5): 146(1987).

72. H. Hanselka and A. S. Henmann, 7.Internationales Techtexil Symposium1994, Frankfurt, June20-22 (1995). 73. M. Avella and R. dell’Erba, Proceeding of the 9th International Conference on CompositeMaterials Vol. 11, Madrid, July 12-16 Vol. 9, p. 864 (1993). 74. A. C. Khazanchi, M. Saxena and T. C. Rao, Textile Composites in Building Construction 1 : 69 (1990). 75. Y. M. Mai, M. I. Hakeem and B. Cotterell, J . Mat. Sci., 18: 2156 (1983). 76. R. S. P. Coutts and A. J. Michell, J . Appl. Polym. Sci.: Appl. Polym. Symp., 37: 829 (1983). 77. R. S. P. Coutts, Composites, 15(2): 139(1984). 78. R. S. P. Coutts, J . Mat. Sci. Letters, 6: 955(1987). 79. R. Selzer, SAMPE, The Materialsand Processes Society, Kaiserslautern, March 28 (1995). 80. M. K. Sridhar and G. Basavarajappa, Indian J . Text. Res., 7(9):87(1982). 81. E. T. N. Bisanda and M. P. Ansell, J . Mat. Sci., 27: 1690 (1992). 82. S. H. Zeronian, J . Appl. Poly. Sci., 4 7 445 (1991). 83. A. N. Shan and S. C. Lakkard, Fiber Sci Techn., 15: 41 (1981). 84. S. C. 0.Ugbolue, Textile Institute 20(4): 1(1990). 85. A. C . Khazanchi, M. Saxena andT. C. Rao, Textile Composites in Building Constructions 69 (1 990). 86. A. J. Michell and D. Willis, Appita, 31f.3):347(1978). 87. H. Saechtling, International Plastics Handbook,Hanser, Miinchen (1987). 88. E. W. Wuppertal, Die Textilen Rohstoffe, Dr. SpohrVerag, Frankfurt (1981). 89. E. Treiber, Die Chemie der Pflanzenzellwand, SpringerVerlag, Berlin (1957). 90. J. Wanvicker, J . Appl. Poly. Sci., 1: 41(1969). 91. S. M. Lee and R. M. Rowell, InternationalEncyclopedia

of Composites, VCH-Publishers Inc., New York

(1991). 92. S. K.Pal, D. Mukhopadhyay, S. K. Sanyal and R. N. Mukhejee, J . Appl. Polym. Sci., 35: 973 (1988). 93. E. P. Plueddemann, Interfaces in Polymer Matrix Composites, Academic Press, New York (1974). 94. M.A. Semsarzadeh and D.Amiri, Polym.Eng. Sci., 25(10): 618 (1985). 95. N. N., ATZ, 97(5):293(1995). 96. G. Wacker and A. K. Bledzki, Kompozyty I Kompoz,ycje Polimerowe, Stettin (Poland), June 22-24 (1994). 97. W.Aichholzer, 14. StuttgarterKunststoff-Kolloquium, Stuttgart. March 22-23 (1995).

53 Short Fiber-Reinforced Plastic Composites Jayamol George and Sabu Thomas Mahatma Gandhi University, Kerala, India

1.

INTRODUCTION

The namepolymerand its applications havenow reached the life style of the masses. Polymer isa generic name with its main permanent members being elastomers, plastics, and fibers. When considering materials for load bearingapplications, designers are increasingly examining the advantages ofusing plastic materials, both thermosets and thermoplastics. The advantages of thermoplastics over others are their low specific gravities and lowcost of fabrication, particularly when using the injection molding technique. The term composite is used in material science to mean a material made up of a matrix substance containing reinforcingagents. By definition, composite is a material composed of two or more distinct components. Composites are divided into twobasic forms, composite materialsand composite structure. Composite materials are composed of reinforcing structures surrounded by a continuous matrix, whereas composite structure exhibits a discontinuous matrix.

B. Fiber-ReinforcedComposites

Fiber-reinforced composites consist of fibers ofhigh strength and modulus embeddeda matrix in withdistinct interfaces between them. In this form, both matrix and fibers retain their physical and chemicalproperties and exhibit synergism. Fibrous composites are divided into two broad areas, continuous fibers and short fibers. Fibers of length less than 50 mm are generally accepted to be short fibers, whereas fiber lengths greater than 50 mm are regarded as continuous. Incorporation of short fibers into polymer matrices offer many attractive features, such as ease of fabrication and better economics at both the incorporation and fabrication stages, i.e., production of a complex-shaped article is possible with short fibers, but iscompletelyimpracticalwithlong fibers. C.

II. CLASSIFICATION OF COMPOSITES

Laminated Composites

In laminated composites, the individual layers of materials are bonded together to form an element or plate. When the constituent materials in each layer are the same, the laminated layers are called a laminate (e.g., plywood, papers, etc.). If the layers are of different constituent materials or of the same material with different reinforcing patterns, the laminate is said to be a hybrid laminate.

The composites have been classified in general under the followingheads: particulate composites, fibrous composites, and laminated composites [ 13.

A.

shape, configuration, or size. Concrete and wood particle boards are two examples of particulate-reinforced composites.

ParticulateComposites

Particulate composites consist of particles dispersed in a matrix. These particles are divided into two classes, skeletal andflakes. The first one consists of continuous skeletal structures filled withone or more additional materials. Flakes consist generally of flat flakes oriented parallel to each other. These particles may have any

D.

HybridComposites

Composites containing morethan one type of fiber are commonly knownas hybrid composites. The term “hy811

812

George and Thomas

1 undlrectional

a M ~ x e dtibrc eq tows

mixed tows

b Mixed hybrid tape ply fibre

2 bidirecticnal eg chop strong mal

q hybrid fabric

C

Individual fibre PlY

e9 ,conventional eq conventional mvlqs

fabrics

d Carbon shell glass core 'sandwich'

Figure 1 Differenthybridconfigurations.

brid" is generally used to denote the incorporation of two different types of materials into one single matrix. The level of mixingcan be either on a small scale (fibers, tows) or on a large scale (layers). The purpose of hybridization isto construct a new materialthat will retain the advantages of its constituents. There are several types of hybrid composites, characterized according to the way in which the constituent materials are arranged [2]. Figure 1 shows the various types of hybrid constructions. These include sandwich hybrids also known as core-shell, in whichone material is sandwiched between two layers of another, interply or laminated where alternate layers of the two (or more) materials are stacked in a regular manner. In intraply, tows of two or more constituent types of fibers are mixed in a regular or random manner, and in intimately mixedhybrids, the constituent fibers are made to mix as much as possible so that no concentration of either type is present in the material.

111.

FIBERS-MATRICES INTERFACE

A. Types of Fibers Used for Reinforcement Fibers are mainly classifiedinto two categories: natural and man-madefibers [3,4]. These two categories are further classified into various types, as shown in Fig. 2. 1. Man-MadeFibers The man-made fibers are classified into two different categories, regenerated fibers and synthetic fibers, depending on the way in which they are prepared. 2. Regenerated Fibers The fibers that are regenerated either from organic or inorganic systems are classified in this category. Polymers fromnatural sources are dissolved andregenerated after passage through a spinneret to form fibers (e.g.,

813

Short Fiber-Reinforced Plastic Composites

FIBRES

L HANHADE

NATURAL

ANIMAL

PLANT

MINERAL

I

LEAF

I

BAST

I

FRUIT

Figure 2 Classification of fiber.

rayon). These fibers are used for reinforcing plastics and rubbers. Glass fibers are the most common ofall reinforcing fibers for polymer matrix composites. The principal ingredient in all glass fibers is naturally occumng silica (SiOz). The two types of glass fibers commonly used in the FRP industry are E and S glass. Other types are C glass, which is used in applications requiring greater corrosion resistance to acids, D glass for electrical application, A glass for making window glass and bottles, and L glass or lead glassfor applications requiring radiation protection. 3. SyntheticFibers Fibers inwhich the basic chemical units havebeen formed by chemicalsynthesis, followed by fiber formation, are called synthetic fibers. Examples include nylon, carbon, boron fibers, organic fibers, ceramic fibers, andmetallic fibers. Amongallcommercially available fibers, Kevlar fibers exhibit high strength and modulus. (Kevlar is a DuPont trademark for poly [p-phenylene diamine terephthalamide].) It is an aromatic polyamide (aramid) in which at least 85% of the

amide linkages are attached directly to two aromatic rings. Electron diffractionand electron microscopy studies revealed that Kevlar 49 fibers exhibit a radially oriented crystalline organization witha uniform distribution of ordered crystalline material throughout the fiber. Properties of some typical fibers are given in Table 1. 4. NaturalFibers Fibers that come fromnatural sources, such as minerals, animals, andplants, are classified as natural fibers [5,6]. Mineral Fibers Fibers in this category are composed of naturally occurring materials. A good example is asbestos. The most commontype is chrysotile, representing more than 95% of world asbestos production. Chemically it is magnesium silicate (Mg6(OH)4 Si205). Today, use of this fiber is limited because long exposure to it may cause bronchial cancer. Animal Fibers Fibers obtained from living organismsare known as animal fibers, e.g., wool, whichis obtained from domestic sheep; silk fiber, which is produced bythe silkworm

Thomas

814

and

George

Table 1 Properties of Kevlar and Other Competitive Materials

KevlarNylon Specific gravity1.44 2758 Tensile strength (kNm-2) 82.74 Tensile modulus (MNm-2) Elongation (%)

1965

1.44 999 5.52 18

Kevlar 49

29 1717 68.95 5

2

SteelE-Glass

1.45 2758 131.0 2.4

7.86

2.55

200 3

Source: Ref. 1.

in makingits cocoon. Silk, in contrast to all other natural fibers, such as cotton, flax, and wool, etc., does not have a cellular structure. In this respect and in the way it is formed, silk closely resembles synthetic fibers. Plant Fibers These fibers are classified into three types depending on the part of the plant from which they are extracted: 1. Bast or stem fibers (jute, mesta, banana) 2. Leaf fibers (sisal, pineapple, screw pine) 3. Fruit fibers (coir, cotton, arecnut)

Extraction methods, amounts, and lengths of various naturalfibers and their physical and chemicalproperties are given in Tables 2 and 3. B. Matrices

The key difference betweenthermoplastic and moretraditional thermoset-based matrices lies in the'behavior of the matrix during processing.The thermoplastic matrix is not required to undergo a cure process to achieve its final mechanical properties. All that is required is melting,shaping,and subsequent solidification. The range of conventional thermoplastic composites and their resultant property advantages depend on the average length of fibers on molding.

The most important matrix material among common polymeric matrices are polyester andepoxy resins. Cured epoxyresins from the prime matrix for high-performance glass, aramid, and carbon fiber composites [7,8], whereas the unsaturated polyester resins are primarily used in industrial applications. The advantages of the former include resistance to water, a variety of chemicals, weathering, and aging. Polyimidesrepresent the largest class of high-temperature polymersin use in composites today [9,10]. They have service temperatures between 250-300°C. Thermosetting resins are crosslinkedwithperoxide to obtain high modulus, strength, and creep resistance. However, they exhibit extreme brittleness. Thermoplastic resins are easier to fabricate than thermosettingresins [l 13, and they can be recycled. PEEK is a semi-crystalline aromatic thermoplastic.Thermoplastic resins such as PMMA have higher fracture energies of about 1 kJ/m2. The typical properties of thermoplastic and thermosets are listed in Table 4. C. Interface

The fiber-reinforced composite materials include three phases: surface of fiber side, the interface between fiber and matrix, andthe interphase. These phases are collectively referred to as the interface [12]. The characteris-

Table 2 Extraction Methods, Amount, and Length of Various Natural Fibers

Fiber

anuallraspador Banana stem of 1.5 wt% Retting/mechanical Coir 8% ofweighs nut (this 0.1 kg) Retting Jute and beatingkhemical 3-4% of stem Rettinddry scratching Linseed 20-25% of dry Straw beatingkhemical and Retting Mesta Same asjute Palmyrahbeating) (by By hand stalk 0.5 kg per Pineapple By hand/decorticator green 2.5-3.5 of wt% leaves Ramie Decorticator bark of 2.5-3.5 wt% ing)/microbial Manual Sisal leaves green 3-4% of retting/decorticator ManuaVretting Sunhemp stalk green 2-4% of Source: Ref. 5.

(mm) 150

300-900 75l500

-

300-600 900-1500 900- 1200 900-1200

-

sites

568-640 413-1627 4-754

0-215

Plastic Short Fiber-Reinforced

815

Table 3 Physical and Mechanical Properties of Some Natural Fibers

Fiber

Initial modulus Ultimate tensile strength (GN/m2)

Elongation at break (MN/m2)

9-22 34-82 7-20 18

Sisal Pineapple Banana Jute Mesta Flax Sunhemp Palmyrah Cotton

-

( W 3-7 0.8-1 1-4 1.3

-

-

2-4 2-4 7-15 6-7

200-400

-

12.5-17.5 0.25-0.40 2-5 0.3-0.5 0.35-0.65 0.18-0.25 12.5-17.5

1-2

780 760

4-6

Flexural modulus (MN/m2)

-

0.03-0.10

Source: Ref. 5.

tics of the interface are dependent on the bonding at the interface, the configuration, the structure around the interface, and the physical and chemical properties of constituents. As a result, the interface has a strong influence on the property of the composite material. These interface problems are seen as a type of adhesion phenomenon and are often interpreted in terms of the surface structure of the bonded material, i.e., surface factors such as wettability, surface free energy, the polar group onthe surface, and surface roughness of the material to be bondedare often discussed as means of improving the bonding strength. A method for the estimation of composite material performance from the characteristics of fillers and the matrices and fromthe configuration of filler is generally called the law of mixture. In the most basic formof the law of mixture, the characteristics of a composite material are represented as a function of characteristics of constituent components and their volume fractions, as shown in Fig. 3. For a composite material (characteristics:Xc) that consists of component A (characteristics: XA ,volume fraction: +A) and component B (characteristics: Xs,volume fraction: + B ) , the basic formulae of the law of mixture are as follows: X, = +AXA+

(1)

+&B

XA

Figure 3 Relation between the properties of composites and various laws of mixture.

The parallel andseries models (curves 1and 2) are given in Fig. 3.

Table 4 Representative Properties of Some Polymeric Matrix Materials Epoxy Polyimide PEEK Polyamideimide Polyetherimide Polysulfone sulfone Polyphenylene Phenolics

35-85 Tensile strength (MPd 15-35 Flexural modulus (MPa) 1.38 Density (g absorption Water 0.10.1 (24 h%)

120

9275

95

105

35

4028

50

35

1.46 0.3

l .30

1.38 0.3

0.25

50-55

70 40

1.25 0.2

1.32 0.2

l .30 0.1-0.2

816

George andThomas

The two curves exhibit theoretical upper and lower limits, respectively, based on a simple composite effect in general. A basic formula that generalizes Eqs. (1) and (2) is:

+AXA" + (3) where n (- 1 d n d 1)represents the properties of the combination mode, i.e., the parallel model is predominant whenn is close to 1 and the series model is predominant when n is close to - 1.

X,"

=

log

=

+A

log XA

+ +B

log x,

(4)

This function is intermediatebetween the parallel model and the series model and referred to as thelogarithmic law of mixture shownin curve 3. The law of mixture is valid for a composite system when there is nointeraction in the interface. However, it is natural to consider that interaction will occur in the interface due to contact between A and B. Then considering the creation of interfacial phaseC , different fromA and B , the following equation can be presented:

~AX+ A +&B + K+A+B (5) This is referredto as quadratic law of mixture shownin curve 4. The parameter K involves an interaction between componentsA and B and provides an expression for the interfacial effect. =

W.

PROCESSINGTECHNIQUES

The choice of manufacturing technologyfor the fabrication of fiber-reinforced plastics or composite materials is intimately related to the performance, economics, and application of the materials. It also depends upon a number of factors, such as component numbers required, item complexity, numberof molded surfaces, and type of reinforcement.

A.

InjectionMolding

One of the most common processing methods for thermoplastics is injection molding. Injection molding offers the advantage of rapid processinginto complex shapes. The list of artifacts manufactured using this process is endless and include electric drill casings, gear wheels, telephones, brief cases, etc. In injection molding, the polymer is fed froma hopper into a heated barrel where it softens and becomes a viscous melt. It is then forced under high pressure into a relatively coldmold cavity where the polymer has sufficient timeto solidify. Thenthe mold is opened andthe fabricated part is ejected. The cycle of operation is then repeated. The processing of discontinuousfiber-filledthermoplastics owes much to therheology ofthe system. Shear thinning andadiabatic heating associated with the broad

molecular weight polymers these in compounds offer attractive processing conditions for engineering applications. During injection molding the shear dominated flow near the surface tends to align the fibers along the flow (shell region), whereas if the flow is extensional in the plane of the part, the fibers near the midplane of the cavity tend to align transverse to the flow (core region) [13-161. It has been experimentally observed by many researchers in the past that the relative extent of coreshell regionsacross the thickness of the injection molded parts strongly depends upon the processing conditions during injection molding andthe properties of the polymer used. Sanon et al. [l31 found that cavity thickness also had anappreciable effect on fiber orientation. During injection molding,a thin layer of solidified polymer is developednear the cold mold walls. The thickness of the cold layer increases near the entrance and diminishes near the advancing front. Thisgrowing-diminishing thermal boundary layer results in a gap-wise converging-diverging flow even in a uniformly thick cavity [17]. This can affectthe fiber orientation in the injection molded parts. The fiber length distribution is determined by the following three stages in processing: 1. The plasticization stages of extrusion compounding 2. Remeltingduringinjectionmolding 3. The higher shear duringinjectionand mold filling

The plasticization stagedominates the fiber attrition process. The flow into the mold cavity during injection dictates the fiber orientation of the final moldedpart and, hence, the properties of the component. The addition of fibers makes polymersystems more shear thinning and thermally conductive. These factors influencethe velocityand shear profiles of the melt front as it passes through the narrow mold channels. This also influences the rate of solidification. The central region is last to solidify and is free to relax into random in-planeorientations constrained by the decreasing channel thickness. This results in skin-core structure. The fiber damage in injection molding severe. is In plunger injection molding, the degradation isalso significant. Increasing the diameter of the nozzle andgates and loweringthe melt viscosity and backpressure reduces the extent of degradation for both thermoplastics and thermosets [18-211.

B. Hand-Lay or Contact Molding Here the mold is coated with a proper release agent to prevent the item stickingon the mold surface. Gel coat resin is appliedby painting or spraying to a thickness of 0.5 mm. This is a pigmental resin that will give the required quality of finish to the component to provide a degree of environmental protection and to prevent the fiber pattern from showing through the external part sur-

817


face. The laminating resin and reinforcement are worked onto the surfaceby roller and brush to the required thickness. The objective of this is to ensure that all the reinforcing fibers are fully coated by the resin. The overall resin-fiber ratio must be controlled such that the mechanical properties of the composite materials are optimized. The item is allowed to cure before being removed from the mold. Curingis camed out at room temperature and can be varied depending on the size and complexity of the item. The cure rate can beadjusted by controlling the resin chemicalformulation, i.e., by varying the catalyst and accelerator levels. The curing process can be accelerated by carrying out a postcure for several hours.

C. SprayMolding Spray molding is the modification of the hand-lay process where the resin and glassfiber are deposited simultaneously onthe molding tool.The fibers are mixed with the resin at the spray head before being deposited on the mold surface. Subsequent consolidation of the laminate is achieved by rolling in a similar manner to the hand-lay process. This method is suitable for large components. Here the capital cost is higher andthe process is very operator sensitive. D. ResinTransferMolding

In this process, resin is injectedinto a closed mold containing the reinforcement preform. The resin can be injected either under pressure [22] or under vacuum [23]. The potential advantages of this process are: (1) low mold cost, (2) inserts can be incorporated, (3) low pressure requirements, (4) accurate fiber orientation, (5) automation possibilities, and(6) versatility. The resin formulation and process variables are selected so that no significant polymerizationoccurs until the mold cavity has been completely filled. This is achieved by the ad-

justment of supply pressure or flow rate mold temperature and catalystlaccelerator levels[24]. The overall cycle time is often limited bythe time required to heat and cure the resin at the injection point.

E. Flexible Resin Transfer Molding Flexible resintransfer molding (FRTM) isan innovative composite manufacturing process developed based on detailed cost analysis [25]. FRTM is a hybrid process that combines the technical characteristics and favorable economics of diaphragm forming and resintransfer molding. Separate sheets of fiber and solid resin are placed between elastomeric diaphragms and heated so that the resin liquefies.The fiber and resinare then compacted by drawing a vacuum between the diaphragms and formedto shape by drawingthe diaphragm assembly over hand tooling (Fig. 4). The process control system combines an empirical resin polymerization model, a fluid flow model, and dielectric sensing of in situ resin properties in order to determine the optimum time for compaction and forming. In the resin polymerization model a sheet of resin is placed between rubber diaphragms along withthe reinforcing fiber material during the FRTM process, the temperature is increased to an isothermal hold at the resin's cure temperature. As the temperature increases, the resinviscosity decreases until the polymerization reaction is activated at the beginning of the isothermal hold. As the polymer molecules begin to grow and molecular weightincreases, the viscosity of fluidincreases. In the polymerization model, the degree of cure a is zero at the time of the viscosity minimum. Thus, the effects of resin aging and resin heatup rate on initial degree of cure are neglected. In the fluidflowmodel,simulation is based on Darcy's law for the steady flow of Newtonian fluids through porous media.This law states that the average

Figure 4 Schematic representation ofFRTM

process.(FromRef. 25.)

818

George and Thomas

fluid velocity U is proportional to the average pressure gradient (VP).

U = -KIP V P

(6)

The constant of proportionality is KIP, where K is the permeability and p is Newtonian viscosity. The dielectric properties of the resin are also measured usingsensors. These measurements werecorrelated with viscosity and used as a part of the FRTM control system. F. ColdPressMolding

This process is carried out at low pressure and at ambient temperature. The fiber is placed in the tool in the form of a mat and the resin is poured in. The mold is closed, and the resin spreads through the reinforcement therebyimpregnatingandfullywetting out the fiber strands. The process is done on an hydraulicpress. Pressure levels requiredfor molding are typically aroundlo’ MPa. Gel coat can be applied to the mold faces before the molding operation commences.

G. Hot PressMolding

at a mixing head and injected into a mold at relatively low pressures. RRIM refers to the composite processing route in which the reinforcement, generally in the form of particulates or very short fibers, is includedin one of the monomer streams. For the most common polymer, the polyurethanes, the polyol would contain the filler. In SRIM,the reinforcement is placedin the mold in the form of a preform. Simple pumps andstatic mixers can be used to disperse the monomers, when the rate of reaction is a lot slower than the fabrication cycle, but effective mixing of fast reacting reagents is best achieved with a selfcreamingimpingementmix-head,which permits the reactants to continuously circulate in their separate systems untilan injection shot is required. The opening and closing of the valve needsto be controlledto accurately deliver the exact quantity of reactants to the mold, otherwise overfilling will cause mold flash and underfilling will cause an imperfect molding. Mixing occurs during mold filling and, more specifically, the dispersion of reinforcement mustoccur effectively. orientation of fibers with higher aspect ratio occurs within the flow fields. For RRIM-based composites, the main benefit is the case with which the properties can be varied across a wide range. Decorative finishes can be applied to RRIM molding by painting or, where color matching between components is less important, by self-coloring. However, color coating of the mold surface prior to RIM provides a better finish.

A premixed sheet or dough material containing resin, fillers, and fibrous reinforcement is inserted between matchedmetaltoolsand the items are compression molded at elevated temperature and high pressure [26]. Different types of molding compounds usedin hot press molding are doughmoldingcompound(DMC), sheet molding compound (SMC), and bulk molding compound 1. Pultrusion (BMC). In DMC the material is suppliedto the molder as a This technique is used to manufacture sectional products dough,which has beenprecompoundedin a mixer. in continuous lengths and is one of the most economiGreat care must be exercised duringthe mixing stageto cally attractive methods for processing thermoset comensure an effective mixing.At the same time, the shear posite materials [30]. Here the fibrous reinforcement is rate should not besevere to minimize the fiber damage. impregnated with thermosetting resin and pulled through In SMC the molding materials are produced in a sheet a heated steel die, often made fromultra high-molecular form where the fiber is sandwiched between two resin weight polyethylene, which shapes the product section layers, and the entire structure is then homogenized by and initiatesthe crosslinking reaction in the resin to form “kneading” betweenspecially countered rolls. This a solid product. sheet is produced continuously and is supplied to the Most pultrusion dies are between 0.4-1 m long. The molder between two thermoplasticcarrier films to preprocessing speed in pultrusion is generally in the range vent adhesion and contamination. BMC is similarin for1-5 dmin”. Some improvement in speed may be obmulation to DMC, whichconsists of thermosetting polytained by preheating the material prior to entering the ester resin, mineral filler, and chopped fiber stands up die. The product is pulled throughthe die by a powerful to 12-mm in length. Pigments, catalysts, etc. are often haul-off unitthat may be either of the reciprocatingtype, included in formulations according to application and where the pultrudate is gripped and pulled in a “hand production requirements. Hydraulic presses are generover hand fashion,” or of the caterpillar type. After the ally used in hot press molding. The molding pressure is haul-off, the product is generally cut by a saw to the in the range of 3.5-15 MPa. required length. Complex sections can be produced by pultrusion. Tubular and hollow shapes require mandrels located insidethe die and anchored upstream of it. Any H. ReinforcedReactionInjectionMolding reinforcement availablein continuous form may be pro(RRIM) cessed by pultrusion. Glassfiber is most commonly used Reaction injection molding (RIM) describes [27-291 the often as alternative layers of unidirectional rovings and process whereby two liquidreactants are rapidly mixed continuous strand mat.

819


V.

FACTORS INFLUENCINGCOMPOSITE PROPERTIES

A.

FiberVolumeFraction

Longitudinal direction

One ofthe important factors affecting composite properties is the amount of fiber it contains, i.e., percentage by volume. Fiber volume fraction (V,) is defined as the ratio of fiber volume (V,) to the total composite volume ( V c ) .

Vf

=

"f v,

Transverse direction ___)

Matrix volume fraction (V,) is defined as:

Using the shear lag theory, the minimum spacing between fibers is determined upon whichthe maximum volume fraction obtained is reported [3 l]. For a composite made of givenfiber and matrixmaterials, there is an optimal spacing betweenfibers at which the fiber tensile strength will be fully exploited. This optimal spacing is the minimum allowable spacing between fibers, below which the structure will start todisintegrate under loading beforethe tensile failure. This minimum spacingthen defines a maximum volumefraction allowable for a composite. B. Strength,Modulus, and Chemical Stability of Fiber and Resin Matrix

The mechanical properties of fiber-reinforced composites very much depend on the strength and modulus of the reinforcing fiber [32-361. Choice of the matrix resin depends on final requirements of the product in addition to low cost, ease of fabrication, environmental conditions, and chemical resistance of the matrix. The function of the resin matrix in a fiber composite will vary, depending on howthe composite is stressed. For compressive loading, the matrix prevents the fibers from buckling and also provides a stress transfer medium, so that when an individual fiber breaks, it does not lose its load carrying capability. The physical properties of the resin influencing the behavior of shrinkage during cure are modulus of elasticity, ultimate elongation, tensile and flexural strength, compression, and fracture toughness. C.

Influence of Fiber Orientation

Orientation of fibers relative to one another has a significant influence on the strength and other properties of fiber-reinforced composites. With respect to orientation three extremes are possible as shown in Fig. 5 . Longitudinally alignedfibrous composites are inherently anisotropic, in that, maximum strength and reinforcement are

Figure 5 Schematic representationof (a) aligned and (b) randomly oriented fiber reinforced composites.

achieved along the direction of fiber alignment. In the transverse direction, fiber reinforcementisvirtually nonexistent. Fracture usually occurs at very lowtensile stress, which may be less than the strength of the matrix. In randomly oriented composites, strength lies between these two extremes. Uniaxial fiber-filledcomposites can have very high longitudinal tensile strength, but the longitudinal compressive strength is generally less because of the buckling of the fiber [37,38]. Transverse compressive strength is limited by the strength of matrix and so is less than the longitudinal compressive strength [38]. By randomly orientingfibers in a plane or by making multilayered laminatesin which the fibers in the various layers have differentorientation directions, composites can be constructed that are essentially isotropic in plane, i.e., such composites have dl irable properties in all directions in a plane. If fibers .e aligned in allthree directions, desirable properties can be achieved in three dimensions. However, to achieve good properties in two or three dimensions, there must be a sacrifice compared with the longitudinal direction of a uniaxially oriented fiber composite. Figures 6a and 6b show the effect of fiber orientation on the tensile strength and modulus of PALF-LDPE composites [39]. D. Influence of Fiber Length

The strength of a fiber-reinforced composite depends not only on the tensile strength of the fibers, but on the

George and Thomas

*

‘

LONGITUDINAL

*

RANDOM

1.400

7-

~

i

LONGITUDINAL

! *

,

1,200

RANDOM

TRANSVERSE _.

..

m

n 1.000 2

m-

3 3

800

n

0 600

! v)

z

2

400

200

-l

0

10

5

(4

15

20

25

30

0

35

0

(b)

FIBRE CONTENT, wt%

5

10

15

20

25

30

35

FIBRE CONTENT, wt%

Figure 6 Effect of fiber orientation on (a) tensile strength and (b) tensile modulus of PALF-LDPE composites.

degree to which an applied load is transmitted to the fibers. The extent of load transmittance is a function of fiber length and the magnitude of the fiber-matrix interfacial bond.The critical aspect ratio that would result in fiber fracture at its midpoint can be expressed as:

Sf (m), = 27

where l

(9)

length of fiber, d = diameter of fiber, (l/d), = critical aspect ratio, S’ = tensile stress at the fiber, and T = fiber-matrix interfacial shear strength. =

The rule of mixtures for discontinuous fiber composites may be expressed as:

where S, = tensile strength of composite, S, = tensile strength of fiber, V, = volume fraction of fiber, V , = volume fraction of matrix, and S, = tensile strength of matrix. More detailed discussion on the interface of fiber length is given in Section X, Theory and Mechanics of Reinforcement.

E. CouplingAgents

Studies on the composite materials have shown that the bonding betweenthe reinforcingfiber and the matrix has a significant effect on the properties of the composite. Good bondingat the interface can be achievedby modifying the fiber-matrix interface with various surface reactive additives or coupling agents. An important technique for improving compatibility and dispersibility between filler and matrix to develop is a hydrophobic coating of a compatible polymer on the surface of the filler before being mixed with the polymer matrix. Generally, coupling agents facilitate the optimum stress transfer at the interface between filler and matrix. The selection of a coupling agentthat can combine both strength and toughness to a considerable degree is important for a composite material. The most common couplingagents are silane-isocyanates-and titanates-based compounds whose chemical composition allows them to react with the surface of fiber. The structural formulae and properties of major silane and titanate coupling agents are given in Tables 5a and 5b. The coatings are applied to the fiber before they are compounded withthe polymer matrix. Coupling action will be explained in Section VII. Peroxide treatment also has significant influenceon improving the interaction between fiber and matrix [39].

Short Fiber-ReinforcedPlastic Composites

821

Table 5a Structures and Physical Properties of Silane Coupling Agents

Molecular Specific gravity atweight 25°C

Structure

Chemical name Vinyltrichlorosilane

161.5

1.26

Vinyltriethoxysilane

10.3

0.93

Vinyltri(P-methoxy-ethoxy)silane

280.4

1.04

236.1

1.07

y-Methacryloxy-propyltrimethoxysilane

248.1

1.04

N-(P-Aminomethyl) y-aminopropyltrimethoxysilane

222.1

1.03

N-(P-Aminomethyl)yaminopropylmethyl-dimethoxysilane

206.1

0.98

y-Chloropropyltri-methoxysilane

198.5

1.08

y-Mercaptopropyltri-methoxysilane

196.1

1.06

y-Aminopropyl-triethoxy silane

221 .o

0.94

y-Glycidoxypropyl-trimethoxysilane

CH2 - CHCH2OCH2CH2CH2Si(OCH3)3

\d

Source: Ref. 40.

F. Influence of Voids

During the incorporation of fibers into the matrix or during manufacture,air or some volatilemay be trapped in the material. The most common cause of voids is the incapability of the resin to displace all the air that had entered within the rising or yarn as it passed through the resin impregnator. The rate at which the reinforcement passes the resin, the viscosity of the resin, the wettability or contact angle betweenthe resin and reinforcement surface, and the mechanical working of the reinforcement will affect the removal of entered air. A high void content usually leads to lower fatigueresistance, greater susceptibility to water diffusion, andincreased variation in mechanical properties. The amount of trapped air or

volatiles that exist in the cured composite can be estimated by comparing the theoretical density with its actual density. VC =

Pc

-P PC

where V, = volume fraction of voids, pc = theoretical density, and p = actual density. The voids represent a degree of unreproducibility that would be desirable to eliminate because they can limit the performance of these materials by providing paths for environmental penetration into and throughthe composite [42].

George

822

Table 5b Structures and Physical Properties of Titanate Coupling Agents

Chemical name

Specific Molecular gravity weight at 25°C

Structure

957

0.94

Isopropyltridodecyl-benzenesurfoniltitanate 9s (2-3252)

1083

1.09

Isopropyl-tri(diocty1-pyrophosphate)-titanate 38s

131 1

1.10

897

0.94

Tetraoctyl-bis-(ditridecylphosphite) titanate 46B

1458

0.94

Tetra(2,2-diallyloxy-methyl-l-butyl)-bis(ditridecylphosphite) titanate55

1794

0.97

Isopropyltriiso-stearoyltitanate

Tetraisopropyl-bis-(dioctylphosphite) titanate 41B

Bis(diocty1pyro-phosphate)-oxyacetate titanate 138 S

1.09

925

91 1

Bis(diocty1pyro-phosphate)-ethylene titanate 238 S

1.08

Source: Ref. 41.

VI. THEORIES OF ADHESION

A.

MechanicalTheory

According to this theory, the adhesive interlocks around the irregularities or pores of the substrate. A rough sur-

face will have a larger Dotential bonding

area than a

dissoives rubber -particles near the polymer 'surface, leaving aporous structure. A metal may then be deposited from solution intothe porous structure providing a mechanical key. Geethamma and coworkers [431 have

823


performed someelectron microscopy studies to support this mechanism. Figures 7a and 7b show the untreated and alkali-treated coir fibers. During alkali treatment, certain holes are produced on the fiber surface after the removal of lignin. These holes will providean anchoring effect andresult in a better interaction between fiber and matrix.

B. AdsorptionTheory The adhesive macromolecules are adsorbed on to the surface of the substrate and are held by various forces of attraction. The adsorption is usually physical, i.e., due to van der Waals forces. However, hydrogen bond-

ing and primary bonding (ionic or covalent) are involved in some cases. If primary bonds are involved, the term chemisorption is used. This theory assumes a definite interface betweenthe adhesive and adherent. Hydrogen bonding is believed to be important in the bonding of tire cords to rubber. There is good evidence that hydrogen bondingisinvolvedin the self-adhesion of coronatreated polyethylene [44,45]. The direct evidence of chemical bondingin adhesion is observed in silane couplingagents. Mablas et al. [46] have providedstrong evidence for chemical bonding between isocyanate coupling agents andnatural fiber-reinforced polyethylene systems. Although the exact nature of the interactions at the interface may be uncertain, the adsorption theory of adhesion is the most widely accepted mechanism.

C. Diffusion Theory The adhesive macromolecules diffuseinto the substrate, thereby eliminating the interface. It requires that the macromoleculesof the adhesive and adherent have sufficient chain mobility andthat they are mutually soluble. Diffusion will also take place when two pieces of the same plastic are heat sealed. Voyutskii [47] provided experimental evidence for this theory based on autohesion experiments, i.e., bonding experiments when the adhesiveand substrate are identical. He studied the bonding of rubbers at elevated temperatures and found that joint strength increases with increasing period of contact, increasing temperature, increasing pressure, and decreasing molecular weight. The strength is also found to be affected by addition of plasticizers and with crosslinking. D.

ElectrostaticTheory

In thistheory, the adhesion is due to electrostatic forces arising fromthe transfer of electrons from one material of an adhesive joint to another. Evidence in support of this theory includes the observation that the parts of a broken adhesivejoint are sometimes charged [48]. It has been shown that peeling forces are often much greater than can be accounted for by van der Waals forces or chemical bonds. VII. A.

Figure 7 SEM of (a) untreated and (b) alkali-treated coir fiber. (From Ref. 43.)

INTERFACE MODIFICATION Surface Modification of Polymers

1. ChemicalTreatment When a polymer is soakedin a heavily oxidative chemical liquid, such as chromic anhydride-tetrachloroethane, chromic acid-acetic acid, and chromic acid-sulfuric acid, and treated under suitable conditions, polar groups are introduced on the polymer surface and the surface characteristics are improved [49,50]. The sur-

824

George and Thomas

face of the polymer is heavily oxidizedby nascent oxygen generated during reaction, i.e., KzC1-207 + 4H2S04 + Cr2(SO4)3 + K2S04 + 4H20 + 3[0]

(12)

The surface of polyolefine isactivated by treating it with the liquid throughthe formation of polar groups such as >C-, -OH, "COOH, and -S03H. Rasmussen et al. [51] qualitatively determined these polar groups in detail. The following mechanismfor the formation of oxygen-containing polar groups has been proposed[52]. R

I

"cH2-C-CH3

Chromic acid

I

H

H I

H I

0 I1

2. CoronaDischargeTreatment Corona discharge treatment results in the formation of high-polarity functionalgroups, such as carbonyl, at the polymer surface. Various mechanisms have been proposed for the improvement of the adhesive properties

283

285 291287

of polyethylene by corona discharge treatment, some attributing itto electric formation [44],others to hydrogen bonding[53]. Figure 8 represents an ESCA analysis of PE treated by corona discharge [54]. The spectrum of CIS,shown in Fig.8 consists of carbon corresponding to functional groups of ethers, alcohols, peroxides, ketones, aldehydes, acids, and esters, which are not detected in untreated material. Chemical shifts related to the bonding conditionsof Cls are + 4.1 eV for R*COOH, +3.9 eV for R*CO+R, +3.2 eV for R*CHO, +3.1 3.3 eV for R-C*O-NHR, + 1.5 eV for R-C*O-R, eV for R*CH20H, + 1.6 eV for RC*H+R, and 0 eV for R"C*HAH-R. By means of these values, peaks in the figure can becorrelated with various functional groups.

+

3. UltravioletIrradiation Some of the advantages of ultraviolet (UV) irradiation on polymer matricesare: (1) reaction occurs at ordinary temperature and pressure, (2) selective reaction is possible, and (3) light energy can be focused on the surface of thematrix. Modification is carried out by introducing functional groups: (1) by applying UV light to oxidize the material surface or alloying the material to contact a gas or sensitizer to cause a photochemical reaction or (2) by alloying UV irradiation graft polymerization to occur at the material surface. Adhesive strength increases with the increasing degree of treatment. To modify a polymer surface, a carboxylic acid group is produced at the surface by introducing double bonds through irradiation in acetylene followed by photo oxidation [S].

289 BINDINGENERGY

(CV)

Figure 8 ESCA spectra of untreated and corona-treated PE surfaces. (From Ref. 54.)


825

In addition, there are many surface modification processes that use triplet sensitizers to permit oxidation reactions. In a typical process, polyisocyanateis applied on a polyolefin together with a sensitizer such as benzophenone and thenirradiated with UV light. As shown in Eq. (15) the sensitizer has an oxidizing effectto produce hydroxyl groupsover the polymer surface. These hydroxyl groups finally react with isocyanate to provide a functional polymer [56,57]. PE

W

7 OH

02

~ II

WC0)3

~

c (15) o

4. Plasma Treatment The plasma utilized for polymer treatment is generally called nonequilibrium low-temperature plasma [59].In low-temperature plasma for polymer treatment, relativelyfew electrons and ions are present in the gas. Here, energy of electrons are in the range of 1-10 eV. This energycauses molecules of gas A to be ionized and excited. As a result radicals and ions are produced. A

+ er+ er-

A A* A* A

A + + 2e A* e Al. A*. A + hv A I-

-

+e

h

re1. The silane coupling agents undergo chemical action withthe surface of inorganic substances to form an SiOM bond (M: Si atom in glass). adsorbed 2. The silane coupling agent is physically on the inorganic surface. 3. The S i - O H group on the glass surface forms a hydrogen bond withthe silanol group derived from the silane coupling agent. a sheathlike structure 4. The silane coupling forms around the glass fibers. 5. A reversible equilibrium reaction takes place between the hydroxyl group onthe surface and the silanol group derived from the silane coupling agent. The mechanism of action of silane coupling agents can be explained terms in of the chemical bondingtheory. ’

0

In some cases adhesive properties can be improved by UV light irradiation alone [58].

Ionization: Excitation: Radical dissociation: Luminescence: Electron addition:

group represents functional groups, which hydrolyzed to give a silanol group (e.g., methoxy andethoxy), and R’ isfor those groups that have an affinity for and display reactivity to the matrix materials. A large number of studies have been reported on the influence of silane coupling agents in polymeric materials reinforced with natural as well as inorganic fibers [61-651. The possible reaction betweensilanecoupling agents and inorganic fillersare [66-691

+

+

(16)

(17) (18) (19) (20)

The activated particles react with polymeric materials so that polymeric radicals are produced on the surface layer of materials. This causes the surface layer to be oxidized, crosslinked, or decomposed. On the other hand, Ass are produced from molecules of the gas and are polymerized, so that the resultant polymers ofA coat the surface of the material. Crosslinking In the presence of an inert gas, such as He orA r , crosslinking can be introduced into the surface layer of materialbyplasma treatment. Hansen andSchonhorn [60] named this “Crosslinking by Activated Species of Inert Gases” (CASING). Asa result, bond strength is enhanced because crosslinking strengthens the surface layer.

As in Fig.9, the hydrolyzable group in a silane coupling agent is first hydrolyzed to give a silanol group, which then undergoes a condensation reaction with the silanol group on the surface of the inorganic material (i.e., glass, silica, etc.) to form covalent bonding between the coupling agent and the material. However, the nonhydrolyzable functional groupin the coupling agent bonds to the matrix through a chemical reaction. Thus, the coupling agentserves to increase the strength of the composite by providing chemical bonding to connect the glass surface and the matrix. The structure of a titanate coupling agent is expressed by the general formulaR-Ti (OXRY)3.These titanate coupling agentsare useful for improving flexibility and processability [70]. Joseph and coworkers [71]

B. SurfaceModification of Fillers

l . Coupling Agents The structure of the silane coupling agent isexpressed where the R 0 by the general formula (RO)3-Si-R‘,

+

Figure 9 Chemicalbondingtheory.

826

George and Thomas

0

0

C-0

C=O

I

I

I

0

I

N "H

L

A A A

i

A

c

A

L

N "H

c

. .. . . .. .. ., . . . . . * .. .. .. - .. .... ... . . .

.-

a

.

.

- cy- cy- cy%&-

*

C h L cy:

LDPE mabix

Figure 10 Schematic representationof the interfacial region of PMPPIC treated fiber and LDPE.

'O

' Ur-"

W

K)rrc 0 5 ''.C * C 410

have reported the use of isocyanates (toluene diisocyanate, poly[methylenel poly[phenyl isocyanate] in natural fiber-reinforcedthermoplastic composites. Urethane derivative of cardanol wastreated with sisalfibers to improve the compatibility betweenfiber and matrix.It was found that PMPPIC treatment has significant influence on the properties of composites, i.e., increased thermal stability, reduced water absorption, etc. The -N=C=O group in PMPPIC is highly reactive with the -OH group of cellulose and, therefore, a urethane linkage is formed.

-N=C=Ot-HO-CellulOse

-

H 0

I II

-N-C-O-CellulOse

(21) A hypothetical chemical structure in the interfacial area of the PMPPIC-treated composite [72] is shown in Fig. 10. The long-chain molecules present in PMPPIC interact with polyethylene leading to van der Waals type of interaction.

VIII. CHARACTERIZATION OF INTERFACE The type of interaction along the interface will exert a great influence on the various properties of the composite materials. Therefore, to improve the performance of a composite material, itis absolutely necessary to characterize the structures of the interface. Some of the methods for analysis of the interface are ESCA, AES, IR-FTIR, SIMS, and SEM, etc. At present, ESCA is widely used in the surface analysis of elements and the qualitativeanalysis of functional groups. Figure 11 shows the ESCA spectrum of polyethylene treated with

392

295

280

540

528

BINDING ENERGY trV)

Figure 11 Relation betweenESCA spectra and Ar+ treatment time of PE. (From Ref. 73.)

Ar plasma for various periods of time [73]. It is clear that there is a spectrum due to an oxygen atomeven for an untreated sample having zero time. If the treating time is prolonged,the combined oxygenatoms increase and the shape of the carbon spectrum changes. Nitrogen atoms do not appear before and after treatment. It is possible to know how carbon is combined with oxygen from the CISspectrum, especially fromthe skirts on the high-binding energy side. Figure 12 shows an analysis by AES spectrum of the interface of a titanium composite material reinforced with W-Sic fiber [74]. The crosssection of a fiber in its center line is analyzed with a beam less than 50 mmin diameter. The results have revealed that a reaction phase of7-12 pm width is formed on the interface between the Sic fiber and the Ti matrix. Figure 13 shows the plot of the interface in the carbon fiber-A1 composite by SIMS [75]. Here measurements by SIMSindicate that the interface reaction phase consists of A l , Na, Mg, and Sn.The wettability of carbon fiber to Na is good, resulting in goodpenetration of Na into the carbon fiber. The penetrating phenomenon is believed to be the result of intergranular diffusion or the formation of an intercalation compound such as CaNa. Carbon fiber is first covered byNa-SN intermetallic compound, andthen by Mg-Sn intermetallic compound. This phenomenon supposed is to be due to the geometric

827


REACTION

ZONE

(C)

A

Figure 13 SIMSplotof Na, Sn, Mg, and AI concentrations as a function of time ingraphite aluminium composite prepared bythe sodium process. (Analysis starts in graphite fiber and proceeds into the aluminium matrix.) (From Ref. 75 .) 0

50

m0

l50

DISTANCE ACROSS THE FBI3 (m)

Figure 12 AES spectra of the W-Sic composite sample. (a) Schematic diagram of the sample (the shaded regions represent the reaction zone). (b) C and 0 line-scan profiles. The maximumPE noise is indicated by anerror bar. (From Ref. 74.)

effect of the irregular surface of carbon fiber or to a decrease of concentration on the interface. The thickness of the interface reaction phase is 0.45-0.55 pm. The infomation obtained from an application of IR spectroscopy to a surface investigation includes the molecular structure, orientation, chemical reaction, conformation, crystallinity, and so on. Figure 14 shows the ATR spectrum of the etched polyethylene surface treated with a chronic acid group [76]. Absorption bands due to surface treatment appear at 3300, 1700, 1260, 1215, and 1050 cm". The band at 3300 cm" represents the absorption due to the hydroxyl group andthat at 1700 cm" is due to the carbonyl group. The bands at 1260, 1215, and l050 cm" are all due to the alkyl sulfonate group.

0

0

0l1

0 II

CH3-0-S-0-CH2-S-0II II

The formation of these polar groups contributes increased adhesion. Observation of disappearing vinyl groups inthe silane coupling agent and of the formation of polystyrene in the silica by FTIR analysis (Fig. 15) have confirmed the occurrence of a reaction between the polymer and the silane coupling agent [77]. The observation of the spectrum for styrene polymerized onthe surface of silane-treated silica andof the difference spectrum of polystyrene adsorbed on the surface of silica have revealed that there are absorption bands of atactic polystyrene at 1602, 1493, 1453, 756, and 698 cm." The absorption bands at 1411 and 1010 cm-] are related to vinyl trimethoxy silane, and C of the difference spectrum is belowthe base line. This indicates that the vinyl groups of silane react with styrene to form a copolymer.

828

,"

1

0

3000

WAVENUMBER

(cm")

J

George and Thomas

Figure 14 Surface IR spectra of etched LDPE-ATR spectra recorded with a KRS-5 reflection element, at 45" angle of incidence. Times refer to chronic acid itch duration. (From Ref. 76.)

AA

Cltio-SLWrrr)Wm

twrsm*Qstx

h

_n^JI

B.

c.

A*

SEM can also be used for the interface analysis of Figures composites. 16a and 16b show the SEM of PMPPIC-treatedand untreated pineapple fiber-reinforcedLDPE composites. Strong adhesion between fiber and matrix is evident from Fig. 16a, whereas Fig. 16b indicates fiber pullout [78]. Figure 17 shows the interface interaction between silane-treated E glass and polystyrene examined by a Raman spectrum [79]. Comparison of B in Fig. 17 with C indicates that the polymerization of styrene is proceeding onthe silane-treated glass, and a comparison of C with D indicates the interaction between the silane coupling agent andstyrene and homopolymerization of the styrene are taking place following the shift of absorption from 1718 cm" to 1702 cm", as carbonyl stretching vibration of the silane has revealed.

L A # - & $ + ""

"""

nu

1

1

1800

1

1

1400

1

1

1

1000

1

1

600 cm"

Figure 15 High-surface area silica treated with aqueous solution of 1 wt%vinyltrimethoxysilane. A silica waspolyIX. INTERFACIAL BOND STRENGTH merized with styrene and washed with CS2 three times. DETERMINATION Polystyrene producedin experiment A was deposited with B silica and the silica washed with CS2 three times. (From Optimum mechanical properties in composite materials 77.) Ref. are strongly related to the efficiency load of transfer.

829


A

B.

C.

D.

I

Figure 17 Raman spectra of a glass fiberhatrix interfaces. (A) styrene monomer; (B) untreated E-glass fiber coated with polystyrene, (C) E-glass fiber treated with 7methacryloxy propyl trimethoxy silane.

(b)

Figure 16 SEM of tensile fracture surface of (a) PMPPIC treated and (b) untreated PALF-LDPE composites.

When an external stress is applied to the specimen parallel to the fiber axis, the tensile load will be transmitted from the matrix to the fiber through the interfacial shear stress. An increase in the applied stress will result in a progressive breakup of the fibers into shorter fragments and a knowledge of the distribution of fragment length gives a value for the critical fiber length and, hence, interfacial shear strength to be determined. On the basis of the Kelly-Tyson model [80] the tensile stress (ax)of the fiber increases linearly from nearlyzero at the fiber ends to a maximum value limited bythe ultimate tensile strengths of fiber (U,-).When this limit is reached, the fiber should break.However, the interfacial shear stress

which has a maximum value at the fiber ends, decreases to nearly zero toward the middle of the fiber. If the external stress is increased, this fiber breakage process should continue until all fiber fragments are less than a critical length, in which case the tensile stress in the fibers cannot reach U -The minimum fiber length for which the tensile stress can reach the ultimate tensile strength of the fiber is termed the critical length (I,). The Kelly-Tyson modelpredicts that the fiber fracture process should result in a distribution of fragment lengths from1,/2 to l,. Often the distribution of fragment length obtained is broader than the ratio 2:l predicted. This has beenattributed to the existence of flaws in the fiber, thereby causing its strength to depend on length [80,81]. Manykindsofinterfacialbond test methods have been proposed [82-851 to study the load transfer including the embedded single fiber tension test, the embedded single fiber compression test, the microde-

(TJ,

830

bond test, the single fiber pullout test, the bead pullout test, the short beam shear test, the transverse tension test, the transverse flexure test, and so on. Among these, the most popularare the fragmentation test, fiber pullout test, and microindenter pushout tests. In the fragmentation test, a single fiber is embedded in a specimen of pure resin, which should have a high strain to failure. The specimen is loaded in tension parallel to the fiber. When the failure strain of the embedded fiber is smaller than the failure strain of the matrix, the fiber will begin to fracture. During the fragmentation test, the strain in the specimen is gradually increased, the strain in the fiber will also increase and reach a new critical fiber strain characteristic of the new fiber fragment length.In this way, the fiber fragments will be fractured again and the fiber is introduced by shear stresses along the fiber surface. Thus, during a fragmentation.test, a saturation fragment length or critical fiber length, (lc) is reached at a certain applied strain, i.e., an equilibrium is reached between interface shear stress and fiber length at a particular fiber fragment length. The fragmentation process can be followed with a light microscope, and the fragment lengthas well as other features (debonding, matrix cracking, yielding)can be monitoredas a function of the applied strain or load. The acoustic emission method can be used if the matrix is not transparent, but does need careful calibration. The interface shear stress can be derived from the observed fragment length if one assumes that its absolute value is constant over the whole fragment length. Based on simple equilibrium of length, one can derive that in each point along the fiber

George and Thomas

correct information concerning the efficiency of load transfer or interfacial shear strength of polymermatrices, i.e., itdoes not take into account matrix viscosity and strain rate. The saturation fragment length is itself a reliable parameter, which for fibers with equalstrength gives a good indicationfor the stress transfer capability of the interface. Interface bond strength and interface friction stress can be derived from the fragment length and the debond strength [ M ] . Absolute valuesof the two interface strengths, i.e., debond strength and frictional stress, can only be obtained when the actual stress is taken intoaccount. Using Eq. (23), one should be aware that the basic assumptions are violated in reality, and the obtained interface strength will lie somewhere between the real interface debond strength and the interface frictional stress.

A. Single Fiber Pull Out Test In the single fiber pull out test (SFPO), a small portion of the fiber is embedded in the bulky matrix and the interfacial strength iscalculated from the peakload when the fiber is pulled out of the composite. In this method [89], a single fiber is taken and partially embedded in a drop of uncured resin placed on a holder. The resin is then cured with the fiber held upright. The holder, with resin and fiber, is held in a grip attached to the crosshead and then pulled out from the resin. The force pulling the fiber out of the resin is balanced byshear stress at the resin-fiber interface holding the fiber in place. The maximum shear stress occurs as the embedded length tends to zero and is given by: where r,, = Fl2rfle and a = (2GiIbrfEf)'"

where df = fiber diameter. If ~i is constant, integration over the fiber length, ( L ) yields:

If we assume that saturation during the fragmentation test just occurs when the maximum fiber stress does not equal the fiber length, the above formula where L = LC and a-, = uddLC), the strength of a filament of length LC,can be used to calculate the interface shear stress. As the interface shear stress is assumed to be constant and characteristics of the fiber-matrix interface, the obtained value is taken to be the interface shear strength [86]. The fabrication of resin block embedded with a single fiber is very time consuming.It is possible to make a specimen in the form of thin tapes using an extruder coupled with a slit die [87]. Because of the viscoelastic interactions in polymers, the fragmentation test may not always yield the

where F is the pull out force, rf the fiber radius, le the embedded length, Gi the shear modulus of interface, b the effective widthof interface, and Ef the tensile modulus of fiber. The embedded lengthhas a significant effect As the embedded length becomes on the value for rmax. smaller, the value of r increases so the widest possible range of embedded lengths should be measured. Depending onthe fiber-matrix system, a wide range of information may be obtained from this test. Four possible regions on a load displacement trace have been identified that may be associated with different interfacial phenomena: anelastic region, a plastic deformation region, a region in which dynamic decoil of the fiber occurs after fiber matrix debonding, and a region where frictional forces predominate and a stick-slip mechanism may occur. In systems where strong stiff fibers have a strong interfacial bond, failure is often catastrophic and only a value for T , , , ~may beobtained. At long embedded lengths, fiber failure may occur within the embedded length resulting in a high apparent pull out load for a short embedded length. This will distort the calculated

831


value for raV.As the bond strength becomes greater, the ability to successfully pull a fiber from the matrix drop without the fiber failing becomes increasingly harder, and problemsmay be encountered when testing systems with very high interfacialshear strengths. In singlefiber pull out from a micro composite (SFPOM) test, another technique can be used for studying the effect of fiber volume fraction on interfacial shear strength that is not obtained from the SFPO test.

For this, failure interfacialshear strength ( 7 ) is obtained by dividing the maximum load P , by interfacial area A . Sometimes the failure occurs by propagation of a crack that starts at the top and travels downward until the interface is completely debonded. In this case, the fracture mechanics analysis using the energy balance approach has been applied[92]in whichP,,, relates to specimen dimensions, elastic constants of fiber and matrix, initial crack length, and interfacial work of fracture ( Wi).

B. MicrodebondTest The single filament pull out test, sometimes called the microdebond test, has receivedattention for some years as a way to assess the adhesion between fibers and matrices in fiber composite [90,91]. It provides a direct measure of interfacial adhesion and can be used with both brittle and ductile matrix resins. Figure 18 shows a widely used test configuration where the matrix isa sphere of resin deposited as a liquid onto the fiber and allowed to solidify. The top end of the fiber is attached to a load-sensing device, and the matrixis contacted byloadpointsaffixed tothe crosshead of a load frame or another tensioning apparatus. When the load pointsare made to move downward, the interface experiences a shear stress that ultimately causes debonding of the fiber from the matrix. The analysis depends on whether the interfacialfailure occurs by yieldingor by crack propagation. The simplest analysis is based on interfacial yieldingwhere the shear stress is assumed to be distributed uniformly over the interface from top to bottom. According to this analysis, the interfacial shear stress increases uniformly until every location in the interface gives way simultaneously.

(25)

where rf = fiber modulus, R = matrix radius, Ef = fiber modulus, E, = matrix modulus, V , = matrix Poisson ratio, Wi = interfacial work of fracture, le = original embedded length of fiber, and a = area of small initial crack at the top of the interface. In the push out test [93],the fiber is pushedinto the matrix rather than being pulled out. The test allows the measurement of two quantities, Fdeb (the force at which debonding occurs) and F p i (the force needed to push the fiber through the matrix sampleif it is thin enough). The bond shear strength 7deb is calculated usingthe shear lag theory:

and

where Ef, E,,, are the elastic moduli of the fiber and the matrix, respectively, and R represents a distance far away in the matrix so the stress or strain at this point is equalto the applied stress or strain. If the testis performed on a thin specimen, then:

t

f

where A x is the sample thickness.

X. THEORYAND MECHANICS OF REINFORCEMENT

Figure 18 Geometry of the microdebond test.

A discontinuous fiber composite is one that contains a relatively short length of fibers dispersed within the matrix. When an external load is applied to thecomposite, the fibers are loaded as a result of stress transfer from the matrix to the fiber across the fiber-matrix interface. The degree of reinforcement that may be attained is a function of fiber fraction (Vf),the fiber orientation distribution, the fiber length distribution, and efficiency of

George and Thomas

832

Figure 19 The representative element usedwith Cox analysis. The miner cylinder represents the fiber and outer as the

matrix.

stress transfer at the interface. In general, the reinforcement is more effective when V, is high, the fibers are long, the fibers are aligned in the principal stress direction, and the interface is strong. There are two well-accepted models for stress transfer. In the Cox model [94] the composite is considered as a pair of concentric cylinders (Fig. 19). The central cylinder represents the fiber and the outer region as the matrix. The ratio of diameters (rlR) is adjusted to the required V,. Both fiber and matrix are assumed to be elastic and the cylindrical bond between them considis ered to be perfect. It is also assumed that there is no stress transfer across the ends of the fiber. If the fiber is much stiffer than the matrix, an axial load applied to the system will tend to induce more strain in the matrix than in the fiber and leads to the development of shear stresses along the cylindrical interface. Cox used the following expression for the tensile stress in the fiber (us)and shear stress at the interface (7): cos h PR,

l,

1

T=Efim[

P=

2Ef In V,-

[2Efr2 In

]

IR

10

Sin h P(R, - X,) (30) coshpR,

l’”

2GmVf-It2

where Ef and G , are tensile modulus of the fiber and shear modulus of the matrix, m is the applied strain, R, is the aspect ratio of the fiber, Ll2r and Xr are the distances from the fiber end measured in terms of fiber diameter. The stress distribution alongthe fiber is shown in Fig. 20. The tensile stress in the fiber rises from zero at the end of the fiber to amaximum at the center. The rate of stress buildup fromthe fiber ends is greater when the parameters G,/Gf and r/R (which is related to volume fraction) are greater. The overall reinforcing e a ciency is givenby the ratio of the area under the curve and the enclosing triangle. This tends to unity as the fiber aspect ratio tends to infinity. The shear stress is greatest at the ends of the fiber and decays to zero somewhere along it. The tensile

7) n n(

””””””””

cf

Gm

”””

Figure 20 The variation of shear stress, T, and the tensile stress along a short fiber in a matrix for short fibers of varying Vd ratio.


833

stress is zero at each end of the fiber and reaches a maximum at the center. If the fiber is just long enough, the maximum tensile stress reaches the tensile stress in the matrix. The ratio (l/d),that occurs under this condition is calledthe critical elastic aspect ratio [94], i.e.: Uf (114, = df

2ri

(32)

For values of lld less than (lld),, the tensile stress in the fiber is alwaysless than that in the matrix. The transfer of load from the matrix to the fiber is poor andthe mechanical properties of the fiber are not fully utilized. If l/d > (l/d),, the tensile stress at the interface remains at a maximum over a greater proportion of fiber length. Here, the transfer of stress from the matrix to the fiber is very efficient, but the average tensile stress in the fiber is always less than that in the matrix because of reduced tensile stress at the end of the fiber. By increasing or decreasing concentration of the coupling agent(l/d),can becontrolled. A coupling agent may increase ri, which in turn decreases the (l/d),. When a fiber breaks, the normal stress at each of its broken ends become zero. Over a distance of 1,12 from each end, stress builds back up to the average value by shear stress transfer at the fiber-matrix interface. Also, the stress state in a region close to the broken ends contain the following: 1. Stress concentration at the void created by the broken fiber. 2. High shear stress concentration in the matrix near the fiber ends. 3. An increase in the average normal stress in adjacent fibers. When a transverse tensile load isapplied, the fibers act as hard inclusions in the matrix instead of load carrying members. Although the matrix modulus isincreased by the presence of fibers, local stresses and strains in the surrounding matrix are higher than the applied stress. The second model introduced by Kelly and Tyson [95] is based on the concept of frictional stress transfer at the interface. It is considered that a constant shear stress is induced from the fiber ends. This results in a linear stress buildup. The frictional stress may be regarded as the interface shear strength (ri).This concept is often used for the experimental estimation of interface shear strength by fragmentation and the pull out test. This model givesa precise definition of the transfer aspect ratio R,, or length. It should be noted that this transfer region increases as the applied strain is increased. If the fiber aspect ratio exceeds 2Rt, there will be a plateau region along the central portion of the fiber. The reinforcement efficiency is obtained from a similar ratio of areas. Most discontinuous fiber composites contain fibers with a wide distribution of lengths distributed in a semi-random orientation, most often close to planar

random. Short fiber systems are more complexthan the models due to fiber length distribution (FLD) and fiber orientation distribution (FOD). If the fiber aspect ratio distribution andthe interfacial shear strength are known, the efficiency of each fiber may be determined using the Kelly-Tyson model [95]. The simplest ,procedure is to use the aspect ratio and orientation data to produce two constants, which may be appliedto the Voigt equation:

E, = qoq1EfVf

+ EmVm

(33)

where T~ and qo are the length andorientation constants, respectively.Theyhave a maximumvalue of unity, when the reinforcement efficiency is equal to that of continuous fibers. The Bowyer and Bader [96] methodology can be used to predict stress-strain response of short fiber-reinforced plastics. The stress on the composite (U,)at a given strain ( E , ) can be computedby fitting the response to a form ofEq. (4) with twoparameters, the fiber orientation factor (C,) and interfacial shear strength (ri).

+ E C E r n ( 1 - Vf 1 where R, = L,ldf - critical aspect ratio Ry R, > R, and L, is the critical fiber length. Since the fiber distribution is larger and R, is a function of E , , RC

C RxVxri n=o is the contribution from the subcritical fiber population and m

C R C

[

1-

z]

the contribution fromthe supercritical fiber population. XI. VARIOUS FIBER-REINFORCED COMPOSITES AND APPLICATIONS

Fiber-reinforced plastics have been widelyaccepted as materials for structural and nonstructural applications in recent years. The main reasons for interest in FRPs for structural applications are their high specific modulus and strength of the reinforcing fibers. Glass, carbon, Kevlar, and boron fibers are commonly used for reinforcement. However, these are very expensive and, therefore, their use is limitedto aerospace applications. Fibrousfillers are nowgainingmore importance over particulate fillers due to the high performance in mechanical properties. The influence of fiber diameter on the tensile behavior of short glass fiber on polyimide was reported [951. At higher concentrations thick fibers seem to be more advantageous probably because of the

834

George and Thomas

higher possibilityfor deformation process in the matrix between fibers with larger separation as in the situation when thick fibers are used. Composite laminates made by reinforcing PF with glass and asbestos find some commercial applications, while those with carbon fiber are being widely evaluated. Current applications include barrier, spacing, and stacking functions inhigh-voltage transformers. The most important is in aircraft applications. Some of the typical fibers and their unidirectional composite properties are given in Table 6. Glass-reinforced nylon is replacing metals in many applications due to their high rigidity, creep resistance, low coefficient of friction, and high heat deflection temperature. Carbonfiber-reinforcednylonhasfounduse in aerospace and tennis racketapplications. The mostwidelyused polyester fiber glass combinations have exceptional strength at low temperature. When combined with specific resins, polyester-reinforced glassfiber can function in temperatures exceeding 400°C. They provide mechanical properties equal to or exceeding those of most metals. Thermoplastic materials are being used increasingly due to their lower costs, light weight, and better flexibility in part design. The use of cellulosic materials in thermoplastic composites is highly beneficial because the strength and toughnessof plastics can be improved.The potential advantages of these natural fiber composites are low cost, light weight, flexibility, reduced wear of processing machinery, and no health hazards. Also, the hollow nature of vegetable fibers impart acoustic insulation properties to certain types of matrices. Different natural fibers are available throughout the world, and these are used as good reinforcements. These composites find immense applications as building materials.The

effects of sisal, coir, bamboo, banana, hemp, and wood 'fibers as reinforcing agents have been studied in an attempt to produce reinforced building products. It was found that pineapple, banana, and bamboo fiber-reinforced cement composites were successful in building applications [97,98]. In 1981, Belmares and hisresearch group in Mexico together with Parfoort and research his group in Belgium carried out an investigation onthe reinforcement of polyester resin with palm, sisal, and henequene [99]. Their work exploredthe reinforcement of polyester resin with palm fibers, a low-cost fiber abundant in Mexico, by improving its mechanical properties, reducing water absorption, andbiodegradation in order to makethem competitive with glass-reinforcedpolyester resin. Palm fibers coated with a poly(viny1 acetate)-fumarate emulsion also increased the flexural properties of composites with polyester resin. Another important aspect of these composites isthe prevention of biodegradation of cellulose fibers in the polyester. When a microbicidal agent (ZnC12) wasadded to the laminated samples, the tensile strength remained almost the same even after 45 days of soil biodegradationin palm fiber-reinforcedpolyester composites [S]. Polyethylene-henequene-sand laminates hasbeen reported to have a relatively lowdensity, high weathering resistance, and low water absorption [loo]. PVC-ixtle-sand laminates [ 1011 show good mechanical properties and are suitable for use as construction materials. In general, building and especially roofingapplications are expected for these composites. Using pinecraft pulp sheets as a basis for polyethylene laminates that show good flexuralproperties at low humidity have been proReduced for application in building and packaging.

Table 6 Comparative Fiber and Unidirectional Composite Properties

Fiberkomposite (GPd

Elastic modulus Tensile strength Density Specific stiffness Specific strength (GP4

(&m3)

(MJW

(MJW

E-glass fiber Epoxy composite

72.4 45

2.4 1.1

2.54 2.1

28.5 21.4

0.95 0.52

S-glass fiber Epoxy composite

85.5 55

4.5 2.0

2.49 2.0

34.3 27.5

1.8 1 .o

Boron fiber Epoxy composite

400 207

3.5 1.6

2.45 2.1

163 99

1.43 0.76

High strength graphite fiber Epoxy composite

253

4.5

1.8

140

2.5

145

2.3

1.6

90.6

1.42

Aramid fiber Epoxy composite

124 80

3.6 2.0

1.44 1.38

86 58

2.5 1.45


835

grafted and ungrafted pulp fibers from aspen and spruce have been combined withpolystyrene by hot pressing, [1021 and as a result, a general improvement of mechanical properties was observed at high-fiber contents. Extruded composites of plasticized PVC and short cellulosefibers have been investigated Goettler by [1031. Pronounced increases in tensile modulus, yield, and ultimate tensilestrength are observed. Single step processing ofreinforcement and polymer with good product performance are key characteristics of the material whose field of application lies inthe vinyl hose industry. The major drawback associated with the use of natural fibers as reinforcement in the thermoplastics matrix to achieve composite material with improved mechanical properties anddimensional stability are the poor wettability and weak interfacial bonding withthe polymer due to inherentlypoor compatibility. Kokta and others [104-1091 have extensively studied the effect of different chemical modifications such as poly(methy1ene) poly(pheny1 isocyanate) (PMPPIC), silane, and monomer grafting on the mechanical properties and dimensional stabilityof natural fiber-reinforced polymer composites (polyethylene, polystyrene, and PVC). They have reported that chemically modified cellulose fiber filled thermoplastic composites offer superior physical and mechanical properties. Felix and Gatenholm [l 101 reported the effect of compatibilizing agent andnature of adhesion in composites of cellulose fibers and polypropylene. Recently, from our laboratory, Thomas and coworkers [ 11 1-1 141 have reported on the use of sisal fiber as a potential reinforcing agent in polyethylene, thermosets (epoxy resin, phenol-formal-dehyde, polyester), and natural rubber. Among various natural fibers pineapple leaf fibers exhibit excellent mechanical properties. George et al. [39,72,78,115]reported the mechanical and viscoelastic properties of pineapple fiber-LDPE composites. Mechanicalproperties of LDPE filled with pineapple fiber are given in Table7. The very close viscosity valuesat high shear rates of these composites for filled and unfilledthermoplastics indicate the successful exploitation of these materials in injection moldingtechnology since very little additional powerwill be required to mold the filled materials.

XII. CONCLUSION At the present time, high strength, high toughness, and long-term durable products can be made from natural fibers and polymers using high technology. By developing low-cost chemicalpretreatment of natural fibers and low-cost manufacturing processes, the integeneous fibers of undeveloped nations will soon be available to produce inexpensive building materials. GLOSSARY AES = auger electron spectroscopy ATR = attenuated total internal reflection BMC = bulk molding compound DMC = dough molding compound e = electron ef = energy of electron Ef = tensile modulus of fiber ESCA = electron spectroscopy for chemical analysis FRTM = flexible resin transfer molding FTIR = Fourier transform infrared spectroscopy Gi = shear modulus IR = infrared spectroscopy K = permeability l/d = aspect ratio LDPE = low-density polyethylene p = Newtonian viscosity VP = pressure gradient PALF = pineapple leaf fiber PEEK = poly(ether ether ketone) PMMA = poly(methy1 methacrylate) PMPPIC = poly(methy1ene) poly(pheny1)isocyanate PVC = poly(viny1 chloride) p = density r- = fiber radius RIM = reaction injection molding RRIM = reinforced reaction injection molding S, = tensile stress of composite SEM = scanning electron microscopy a - = tensile stress Sf = tensile stress of fiber SIMS = secondary ion mass spectroscopy

Table 7 Mechanical Roperties of PALF-LDPE Composites

ensile Fiber break atmodulus strength content (W%)

Hardness set

( M W( M W

130LDPE 61010 20720-900 30 1095-1100 Source: Ref. 39.

8.5 16.3 19.8 22.5

(%)

63

strength

(kN/m)

(%)

Density (&m')

72 81 97

45 55 60 65

166 2 2 1

0.90 0.95 0.99 1.03

110 11

9 4

836

Sm = tensile stress of matrix SMC = sheet molding compound SRIM = structural reaction injection molding 7 = shear strength V, = volume fraction of composite Vf = volume fraction of fiber V, = volume fraction of matrix Wi = interfacial work of fracture

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54 Peculiarities of the Fine Structureof PET Fibers and the Relationship to Their Basic Physical Properties Grzegorz Urbanczyk and Andrzej Jeziorny Technical University of Lodt, Lodt, Poland

1.

INTRODUCTION

a large scale. As for the latter, some unfavorable properties of standard fibers were minimized and some new favorable ones were selectively obtained. Fibers less susceptible to pilling, with improved dyeability, highly shrinkable fibers, or capillary porous fibers of increased hygroscopicity are a few examples.The favorable functional qualities of PET fibers result from their specific physical properties (i.e., mechanical,thermal, frictional, and electrical) and are particularly stronglyconditioned by the specificity of their fine structure. This fact, as well as the fact that the fine structure of these fibers can be shaped during the manufacturingprocess, over a wide range justify the consideration of the basic aspects of their fine structure.

Polyethylene terephthalate (PET) fibers and filaments are currently the most important kind of man-made fibers. The manufacture of these fibers on an industrial scale, initiated in England in 1947and then in the United States, originated with J. R. Whinfield and J.T. Dickson’s patent (British Patent 578079 submitted in Great Britain on July 29,1941). The PET fiber production, which in the 1940s and 1950s fell in the shadow of the then dominating (in regards to the rate of production) polyamide fibers, increased dynamically in the course of time. This dynamic increase was confirmed by the proportion indicesof PET fibers in the world production of synthetic fibers: 34% in 1970,47% in 1980, and 54% in 1990. At the beginning of the 1970s the volume of PET fibers production was greater than the quantity of polyamide fibers produced, which placed them in first position not only among synthetic fibers but among all chemical fibers. In 1994, the proportion of PET fibers in the world production of synthetic fibers was 62.9% and of chemical fibers was 55.3%, while in the total volume of all kinds of fibers it was 27.4%. Out of PET fibers presently produced, 38% are staple fibers and 52.5% are filament yams, with a marked tendency toward an increase in the latter. A 55% proportion is anticipated in the year 2000. At present, about 75% of PET fibers are used for textile purposes and 25% for nontextile purposes. This dynamic increase in production was accompaniedby the qualitative development of PET fibers, which manifested itself in the widening ofassortment of the fibers being produced (e.g., staple microfibers and filament yarns of the POY, MOY,FOY, and HOY type) and in the manufacture of second-generation fibers on

II. FINE STRUCTURE OF PET FIBERS

PET fibers in final form are semi-crystalline polymeric objects of an axial orientation of structural elements, characterized by the rotational symmetry of their location in relation to the geometrical axis of the fiber. The semi-crystalline character manifests itself in the occurrence of three qualitatively different polymericphases: crystalline phase, intermediate phase (the so-called mesophase), and amorphous phase. When considering the fine structure, attention should be paid to its three fundamental aspects: morphological structure, in other words, super- or suprastructure; microstructure; and preferred orientation. A.

Superstructure

PET fibers and filamentsare characterized by a fibrillar superstructure that corresponds to the general concept of the fibrillar structure of synthetic fibers. The fibrillar 839

840

structure concept was established by Peterlin [1,2], developed further by Brestkin [3], Gojchman [4], Pechhold [5], and finalizedby Prevorsek and Kwon[6]. According to this concept, the fiber in macroscopic scale is formed by a system of interconnected fibrils, sometimesreferred to as macrofibrils. Particular fibrils are connected to one another to form a coherent whole by means of tie molecules passing from one to another fibril. The structure of the fibrils is determined by their substructure. Most often this aismicrofibrillar substructure, also known as longitudinal substructure. In the case of this type of substructure, the fibril is made up of a bundle of (3-7) smaller fibrillar elements laidin parallel, socalled microfibrils. The connections betweenmicrofibrils are ensured by tie molecules passing from one to another microfibril. The microfibril, as in the fibril, is an elongated object of a lateral dimension of 25-80 A, and a longitudinal dimension of 1000-10,000 A.Along the microfibrils there occur-regularly, periodically, and alternately-crystalline blocks, also called crystallites, and noncrystalline areas between them, referred to as separating layers. The longitudinal periodic regularityin the structure of microfibrils manifests itselfin the fact that in each fragment of its length, the sum of the height of the crystalline region and the separating layer are constant and equal to the value of the long period or mean long spacing (L). For PET fibers, this long period is within the range of 90-170 A. In the long period, about 50- 100A falls to the height ofthe crystallineblock, while 40-70 A falls to the thickness of the separating layer. The principle of structure of microfibrils is shown in Fig. 1. The periodic regularity of the microfibril structure causes its degree of crystallinity to be constant, and when two-thirdsof the value of L falls to the crystalline layer thickness, it amounts to 0.66 (volume degree of crystallinity) or 0.75-0.8 (mass degree of crystallinity). Crystallitesin microfibrils of PET fibers are created in their central part of the set of parallel, straightened, and spatially ordered segments of molecules to form a crystalline region. The developed space lattice of the crystalline region characterized is by a triclinic crystallographic system. On both ends the crystalline region adjoins so-called crystallite surface layers. The surface layers are formed mainly by regular and irregular chain folds. The length of the regular folds, also called adjacent reentry folds, is much smaller than that of the irregular folds and also the smallest possible fold. The surface layers are characterized by relatively high ordering of the regular chain folds. As shown by one of the coauthors [7], this ordering can be confirmedtheby relatively low value of the surface free energy of these layers. Noncrystalline separating layers of microfibrils are formed from two crystallite surface layers that are side by side in the microfibril and from a central part called the intermediate zones. The latter are mainly made up of tie molecules among whichare distinct straightened

Urbahczyk and Jeziorny

Figure 1 The structureof a microfibril. C - crystallite; S separating layer;SL - surface layer;IZ - intermediate zone; mF - border of microfibril;1, - crystallite length;1, - separating layer length; L - mean long period (spacing).

and stressed molecules(so-called taut tie molecules [TTM]) joiningadjacent crystalline blocks in the microfibril. The intermediate zones also comprise the ends of macromolecule chains, so-called cilia, and coiled ends of molecules. Microfibrils are formed in PET fibers during the stretching of a solidified polymer stream. With an increase in the draw ratio, microfibrils become increasingly slender. The microfibril length is assumed to be proportional to the draw ratio (R), their lateral dimension proportional to while the length-to-lateral dimension ratio is proportional to R3'*. Another type of fibril substructure in PET fibers, besides the microfibrillar type already discussed, is the lamellar substructure, also referred to as the lateral substructure. The basic structural unit of this kind of substructure is the crystalline lamella. Formation of crystallinelamellaeis a result of lateral adjustment of crystalline blocks occurring in neighboring microfibrils on the same level.Particular lamellae are placed laterally in relation to the axis of the fibrils, which explains the name-lateral substructure. The principle of the lamellar substructure is shown in Fig. 2. The distinction between the types of substructures of the microfibrils occumng in the PET fiber is possible based on the value of the substructure parameter ( A ) .

a,

841

Fine Structure of PET Fibers

B. Microstructure

c1

Figure 2 The lamellar substructure of a fibril. (a) Reciprocal positions of crystalline lamellae asa result of fiber annealing. (b) The situation after relaxation of stress affecting TTM. a1.2- average angleof orientation of TTM;CL - crystalline lamellae;CB - crystalline blocks (crystallites); mF border of microfibrils; and F - fibril. In order tosimplify it was assumed that: (1) there are the taut tie molecules (TTM) only in the separating layers, (2) the axis of the fibril is parallel to the fiber axis.

(For details see Section 111. A., Mechanical Properties.) For values of this parameter less than 1, we canconclude that a substructure of the microfibrillar type occurs, whereas for the values equalor close to 1 a substructure of the lamellar-type occurs. While preserving its essence, the superstructure of PET fibers can besubject to modification, dependingon the manufacturing conditions of the fiber, and to additional refiningtreatment. This modificationfirst changes the value of the long period( L ) .The change in the value of L is due to an increase in the height of the crystalline blocks and reduces the thickness of the separating layers. The values, listed in Table 1, are an illustration of such changes.

Characteristic of the microstructure of PET fibers in their final production form is the occurrence of three types of polymer phases: Crystalline, mesomorphous, and amorphous. The first phase is the result of crystalline aggregation of PET molecules, the second phase-of mesomorphous or, in other words, paracrystalline aggregation, the third phase-of amorphous aggregation. The mesomorphous and amorphous phases together form a noncrystalline part of the fiber. l . Crystalline Phase The crystalline phase is formed by a set of crystalline regions that occur in the fiber. In particular, crystalline regions Laveto do with a spatial regularsystem of fragments of molecules that form a crystalline space lattice. In the case of PET fibers, it is determined bya triclinic crystallographicsystem. Developing a crystalline phase is tantamount to the formation of crystallites, which is the result of the process of the fiber polymer crystallization. In the case of PET fibers, stress-induced crystallization, which occurs in the course of stretching the solidified fiber, is decisiveas is thermalcrystallization, which is the thermal treatment of the fiber while manufacturing or refining it. In order for crystallization to occur there must be a sufficiently high value of elongation imparted while stretching and a heat input temperature above the glass temperature (TG). The process of stress-induced crystallization starts at a threshold value of elongation of about 250-300%. The uncoiling, straightening, and parallelization of molecules, occurring during the stretching of a freshly solidified fiber, is accompanied by the conformation of the transformation of the dioxymethylene fragment of the chains. This transformation consists of the transition of the initial stereoisomer of the gauche-type of this group into the stereoisomer of the trans-type (Fig. 3). The quantitative effect of crystallization also depends on the average molecular weight

Table 1 Values of the Long Period of Differently

Drawn PET Fibers

Long period L (nm) (200°C, Before

After annealing Lmin,

bath)silicon Draw annealing ratio 3.014.0 x x 3.5 16.0 4.0 17.0 x 4.5 x 5.2 x

11.2 12.0 13.4 14.0 16.6

18.0 19.0

TRANS

GAUCHE

Figure 3 Trans and gauche isomers of PET.


842

of the polymer used for fiber production. When the remaining productionparameters are on the same level, a smaller average molecularweightusually leads, to stronger crystallization of the fiber. A univocal confirmationof the developmentof crystalline aggregationjn the fiber is the occurrence of layer reflexes Oi 1, 1 i 1 , 1 1 1 ,and 101 on the textural x-ray diffraction pattern. The details of organizationof the space lattice are defined bythe parameters of the unit cell and the number of polymers falling into one cell. The data, established by different authors, are presented in Table 2. Daubenny and Bunn's[8] pioneer findingsare considered the most probable for space lattices occurring in PET fibers. A specific attribute of unit cell buildingthe isinclination of the axis of macromolecule chains, in relation to the normal, to the plane of the base of the cell (ab). According to Yamashita [l11 this inclination is within the range of 25-35" (Fig. 4). Against the background of space lattices of other types of fibers, the lattice of crystalline regionsin PET fibers is characterized by a number of specific features. These are: The perfection of the geometrical structure. This is confirmed by a great number of x-ray reflexes, including, andin particular, the occurrence of layer reflexes and reflexes of a higher order than the first one (Fig. 5 and Table 3). The translational direction of the lattice (c), which is the direction that the crystallite axis is not, because of the triclinic crystallographic system, is perpendicular to the plane of the unit cell base (ab). The space lattice does not undergo polymorphous transformation. As with other kinds of fibers, no transformation of the space lattice under the effect of any physical or chemical treatment of PET fibers has yet been found. Crystallites occurring in PET fibers can assume two kinds of morphologicalforms. The first form represents the crystallite formed by molecules of folded conformation, while the other is formed from molecules of extended-chain conformation.The first form is sometimes called a flexural morphologicalform, whereas the other is calleda straightened morphologicalform. The flexural form is the typical and prevailing morphological form in PET fibers. However, it should be stressed that no

Figure 4 The scheme of the unit cell of PET.

occurrence of the spherolitic morphologicalform of crystalline aggregationhas been observed in PET fibers. Recognition of the kind of a morphological formof the crystallite can be achieved by means of x-ray, IR spectroscopic, or thermal analysis. In the first case, it consists of studying the low-angle scattering of x-radiation. The occurrence of long periodreflections is the evidence of the occurrence of the first morphological form.In the case of IR spectroscopic analysis [12,13], the presence of an absorption band of 988 cm" confirms the occurrence of this form. According Baranova to et al. [l41 and Pakszwer et al. [15], the occurrence of the 853-cm" band can be consideredthe criterion for the presence of crystallites of the first form, while a 846-cm" band points to the second form.

Table 2 Parameters of the Unit Cell of PET

4.56 4.52 4.48

5.94 5.98 5.85

10.75 10.77 10.75

98.5 101.0 99.5

118 118 118.4

112 111 111.2

1.455 1.476 1.515

181

191 U01

pe


843

treatment. The decisive factors are thedraw ratio of the fiber, and the temperature and the duration of action of heat. Similarly, the average size of crystallites is also diversified, andas with the degree of crystallinity, it depends onproduction conditions and the additional refining treatment. The numerical data, as determined by the authors, characterizing the dependence of the degree of crystallinity and the size of crystallites on the draw ratio of the fiber, are presented in Table 4.

2. Mesomorphous Phase

Figure 5 X-ray reflexes of PET fibers differentiated by draw ratio (in %).

The quantitative proportion of crystalline aggregation in the total polymer mass of the fiber, i.e. degree of crystalliniry, is strongly diversified for PET fibers. The degree of crystallinity is within the range of the values 0.10-0.70. Most often, however, itiswithin the range of 0.45-0.65. The degree of crystallinity depends onproduction conditions and the additionalrefining

The mesomorphous phase, also called an intermediate phase or a mesophase, is formed by molecules occurring in surface layers of the crystallites. It can be assumed that the mesophase is made up largely by regularly adjacent reentry folds. However, it cannot be excluded that the mesophase isalso composed of some irregular chain folds, which are characterized by a long length and run near the crystal face in the direction perpendicular to the microfibril axis. The mesomorphousphase is characterized by a paracrystallinearrangementof segments of molecules, analogous to the one occurring in nematic liquid crystals. The paracrystalline manner of the arrangement of the chains manifests itselfin the development of a pseudolattice. Accordingto Asano andSet0 [ 161, this lattice for PET fibers is described by a monoclinic crystallographic system anda unit pseudo-cellof this lattice of the parameters: a = 4.3 A, b = 9.2 A, c = 10.5 A, and (Y = looo. The density of such a pseudo-lattice is 1.376 - 1.392 g/ cm3. The occurrence of the mesophase in the fiber is confirmed byx-ray diffraction examination.The occurrence of three equatorial reflections 010, 110, and 100, the absence of layer and meridionalreflections, and the manifestation of the intensity maximum of diffusively scattered radiation at 20 = 19O in the fiber diffraction pattern are the criterion for the presence of the mesophase. The

Table 3 X-Ray Diffractions in Texture Diffractograms of PET Fibers

Diffraction indices Interplanar spacing Reflex Equatorial Equatorial 3.94 Equatorial 3.48 Layer Refl. Layer Refl. Layer Refl. Layer Refl.2.72 Layer Refl.3.90 Meridional3.58 Meridional

strong strong very strong strong strong medium weak weak very weak medium

17.50

21.43 27.77

( W

(A)

010

5.02

111

5.46 4.14 3.21

110 100 0116.20 1 111 101 112 103 105

2.15

dhkl

Glancing angle for CuId-2Qhkl

("1

41.10


844

Table 4 Degree of Crystallinity and Average Crystallite Size of Differently Drawn PET Fibers

Degree of crystallinity coefficient method Density method X-ray Draw ratio (Xa) 2.0 x 3.0 x 3.5 x 4.0 x 4.5 x 5.2 x

0.12

0.22 0.27 0.31 0.33 0.35

Lattice disorder

(X) 4.8 4.6 3.7

-

0.26 0.29 0.3 1 0.32 0.40

(k) ~

Average crystallite size perpendicular to the crystallographic plane (hkl) Dhrl (nm) (010)

(100)

3.9 3.7

3.2 3.2 3.2 3.0 2.9 2.8

~~

16.5 2.8 2.4 3.1 3.3 4.5

3.6

Ruland's 1421 and Vonk's 1431 methods.

broadening of the base of the curve of the azimuthal intensity distribution of the reflected radiation J = f(6) for reflexes 010, 110, and 100 is an additional confirmation ofthe occurrence of the mesophase. The x-ray identification of the presence of the mesomorphous phase is, however, hindered by the simultaneous occurrence of a crystalline phase with it. The latter is a source of equatorial reflexes that partly overlap the equatorial reflexes fromthe mesophase and dimthe presence of the latter ones. The quantitative proportion of the mesomorphous phase in PET fibers is strongly diversifiedand depends on the conditions of fiber manufacture. According to Lindner [17], it isestimated to bein the range 0.21-0.36. For poorly stretched fibers (below 300% of draw ratio), the mesophase constitutes almostexclusively an ordered part of the polymer in the fiber. With an increase in the draw ratio of the fiber, the proportion of the mesophase in the whole fiber mass increases, mainly at the cost of the amorphous phase. It is assumedthat themesophase constitutes about 0.5 of the amorphous phase mass for poorly stretched fibers and about 1.2 of this mass for highly stretched fibers. 3. Amorphous Phase

The amorphousphase is formed byfiber areas in which the chains of macromolecules or their fragments are characterizedby a very low degree of their spatial ordering. The amorphous phase is first formed by taut tie molecules (TTM)in the separating layers of microfibrils and placed between crystalline lamellae in the fibrils. Tie moleculesoccumng between microfibrils and fibrils are another component of the amorphous phase. The amorphous phase is also made upof irregular chain folds occurring in crystallite surface layers and of cilia. It is assumed that, depending on the degree of chain ordering, the amorphous phase is strongly diversified with respect to the molecular cohesion. The amorphous

phase regionsof the smallest cohesion will be characterized by the absence of intermolecular bonds. This fact justifies referring to polymer materialof such regions as nonbonded amorphous material. In turn, for regions of a greater cohesion the intermolecular bonds will occur, which is the case for bonded amorphous material. The amorphous phase differs from the mesophase and the crystalline phase by a clearly lower value of density. The amorphous phase density depends on the internal orientation of the fiber. Its value isin the range 1.335-1.357 g/cm3. Inthe case of a very highorientation, it can even reach the value 1.363 g/cm3. The existence of the amorphous phase of the fiber is confirmed in x-ray examination by the occurrence of a distinct intensity maximum of the radiation scattered diffusively at 2Q = 21.6'. The fraction of the amorphous phase in the fiber depends on manufacturingconditions and a possible further refining treatment. It is estimated to vary from 0.25 to 0.60. With an increase of the draw ratio and following the thermal treatment of the fiber, the proportion of the amorphous phase only reaches the lower values of this interval.

C. Orientation The development of the internal orientation in formation in the fiber of a specific directional system, arranged relative to the fiber axis, of structural elements takes place as a result of fiber stretching in the production process. The orientation system of structural elements being formed ischaracterized by a rotational symmetry of the spatial location of structural elements in relation to the fiber axis. Depending on the type of structural elementsbeing taken into account, we can speak of crystalline, amorphous, or overall orientation. The first case has to do with the orientation of crystallites, the second-with the orientation of segments of molecules occurring in the noncrystalline material, and the third-with all kinds of structural constitutive elements.


845

1. CrystallineOrientation The parallelization of crystallites, occurring as a result of fiber drawing, whichconsists in assuming by crystallite axes-positions more or less mutually parallel, leads to the development of texfure within the fiber. In the case of PET fibers, this is a specific texture, different from that of other kinds of chemical fibers. It is called axialtilted texture. The occurrence of such a texture is proved by the displacement of x-ray reflexes of paratropic lattice planes in relation to the equator of the texture diffractogram andby the deviation fromthe rectilinear arrangement of obliquediffraction planes. With the preservation of the principle of rotational symmetry, the inclination of all the crystallites axes in relation to the fiber axis is a characteristic of such a type of texture. The angle formed by the axes of particular crystallites (the translation direction of space lattice [c]) and the

fiber axis is called a tilting or inclination angle of the texture (p),(Fig. 6). The direct cause for the formation of the axial-tilted texture in PET fibers is the triclinic crystallographic systemof the space lattice and the specific arrangement of chain molecules within the space lattice (Fig. 4). The value of the angle of tilting of the texture can be determined from the analysis of x-ray fiber diagrams. As Urbariczyk noted [18], the position of layer reflexes Oil and i 1 1 or the position of equatorial reflexes 010 and 100 can be analyzed. In the first case, the tilting angle of the texture (9) can be determined from the equation:

a X cos2(cp) + b x cos(cp)

+c =0

(1)

where:

a = 351.9, b = -451.2 x sinfill - 512.8 X sinsoil, and c = 291.8 X Sin28i11 336.1 X sin2&1 + 435.1 X sin2fi11x sinsoil x 132.8

+

where fil 1 and are azimuthal angles determining the position measured from _the equator intensity maximum of the reflexes Oil and 1 1 1.

crg(a) =

In the second case, from the equation: cos(0)lm x sin(p) = sin(a) (2) where a is an auxiliary angle definedthe byrelationship:

cos(0)oIo x sin(b)olo - cos(59.5") x cos(19)~~ x sin(3)lm sin(59.5") X cos(8)lm X sin(6)lm

where 8100and Solo are azimuthal angles determining the position of the maximum intensity of the reflexes 100 and 010 measured from the equator. It is more convenient to use the procedure of the first kind for higher drawn fibers (over 350% of elongation) and the second procedure for fibers of smaller drawing. The value ofthe angle of inclination of the texture depends on the magnitude of the drawing applied. As a result of the investigations of one of the authors

I

[B], it is known that with an increase in the drawing, the value of the angle p causes the axial tilted in the PET fiber's texture to be similar to the ordinary axial fibrous texture occurring in the majority of chemical fibers. As a consequence, the fiber diagram of PET fibers is practically indistinguishable from the kinds of texture diffractograms of other chemical fibers. The ordering of crystallites in the fiber texture is best and univocally described the in quantitative manner

II

Figure 6 Schematic presentationof the ideal axial-fibrous texture(I) and ideal axial-tilted texture (11).


846

by the value of Hermans’ axial orientation function VC) determined by the dependences:

where 4 is the angle formedby the crystallite axis (the translational direction of the space lattice [c]) and the geometrical axis of the fiber, i.e., called crystallite axis orientation angle. As a rule, to determine the value offc, x-ray diffraction examination isused, which consists in the quantitative analysis of the azimuthal intensity distributions of the x-ray reflexesof the texture diffractogram. Twoprocedures are widely used. Oqe of them consists of considering the meridional reflex 105, the other in the analysis of the equatorial reflexes 010, 110, and 100. In the first procedure, which is considered theoretically to be more proper, the value of fc is generally determined by the dependence in Eq. (4). In this case, the angle @ value depends on the angle parameter of crystallographic(i05) planes ( E ~ O S (the ) angle betweenthe “c” axis of a crystalliteand a normalwhenconsidering crystallographic planes) and on the directional angle of the (105) planes ( ~ ( 7 0 5 ) ) . The equation connecting the mentioned angles is as follows:

4 = €105 - - pios

(5)

Using Eq. ( 3 , the expression in Eq. (4) can be written in the form: fc=lThe correct determination offc depends largely on using the correct value of the €105 angle. In lightof the analysis of a coauthor [191, the value of €105 is equalto €105 35.8O. Assuming this €105 value and the occurrence of the unit cell proposed by Daubeny and Bunn [8], after calculation of values of trigonometric functions, the expression in Eq. (6) may be written in the form [19]: fc

= 0.487 X

- 0.474 pi05

X

X

s i n 2 piOS

(1 -

5

2

+ 1.424 pi’s)

(7)

where sin2pTo5 is the variation in the parameter of the orientation of crystallographic (105) planes derived by Wilchinsky [20]. When using the procedure with equatorial reflections, the value offc is determined fromthe relationship deduced by Gupta and Kumar [22]: fc

= 1

-

- 0.534

X Cosz polo - 1.150 X 1.315 X c o S 2 p 1 ~

Cos2

p170

(8)

where C O S ~ ~ O I COS2plTo, O, and c o s 2 p l are ~ the variation parameters of the orientation of crystallographic (OlO), (1 lo), and (100) planes.

The quantitative assessment of the degree of crystallite orientation by x-ray examinationis not free of ambiguity. Froma comparative analysis [23] in whichresults obtained fromthe consideration of (105) and fromthree different variations of equatorial reflection were compared, the conclusion was that the first procedure can lead to underrated results, i.e., to the underestimation of the orientation. However, it can be assumed that this does not result,from an incorrect procedure, but from ignoring the fact that the adjacent (105) reflex canoverlap. The absence of the plate effect of the orientation is characteristic of the orientation of crystallites in PET fibers. The evidence of this absence is the nearly identical azimuthal intensitydistributions of the diffracted radiation in the reflexes originating from different families of lattice planes. The lack of the plate effect of orientation in the case of PET fiber stretching has to do with the rod mechanism of the crystallite orientation. The orientation of crystallites in PET fibers can also be assessed quantitativelybymeans ofIR spectrographic examination. In this case, the basis for the assessment are the values of dichroic ratio (R) of the “crystalline” absorption bands in the fiber spectrogram. The determination of the values of fc is made usingFraser’s dependence [24,25] modified by Chranowski [26]:

2 x (1 - R ) x cos2 ff - 1)

f = (1 + 2R) x (3

where a is the angle of transition moment for the absorption group correlated with the absorption band taken into consideration. In an investigation [27], the determination was based on the consideration of crystalline bands 875 cm“, 1343 cm”, and 1473 cm”. In the case of absorption groups corresponding to these bands, i.e., C&, CH2, and CH2,respectively, the values of the angles of the transition moment(a)are equal to ffS75 = 71.0°, a1343 = 43.0’, and ff1473 = 43.5’. The crystallite orientation in PET fibers depends first and basically onthe applicated draw ratio and second on the stretching rate. The values offc characteristics for PET fibers as established by the authors are in Table 5 .

2. Amorphous Orientation The amorphous orientation is understood as anarrangement of kinetic segments of macromolecules that form the amorphous phase in relation to thegeometrical axis of the fiber. The development of an amorphous orientation in PET fibers initiated at the stage of melt spinning is intensified in the process of stretching the solidified fiber. Stretching in the production process is a decisive factor in the amorphous orientation in its final form. Developing the amorphous orientation starts at the beginning of stretching, then increases rapidly to reach the draw ratio of about 3.5 x. For a higher draw ratio, i.e.,

Fine

of PET Fibers

847

Table 5 Crystallites Orientation Function(fc) of Differently Drawn PET Fibers ~~

IR

Draw ratio

method

Equatorial reflexes (OlO), (100) (110), ~

2.0 x 3.0 x 3.5 x 4.0 x 4.5 x 5.2 x

X-ray

~~

method spectroscopy Meridional reflex (105)

-

-

0.817 0.820 0.831 0.713 0.842 0.844

0.912 0.916 0.930 0.945 0.932

greater than 4.5 x, an effect of an apparent decrease in the orientation index value can be observed [28]. Such an effect isconsistent with the fact that athighest draw ratios part of the amorphous phase is transformed into the crystallinephase as a result of stress-inducedcrystallization. Consequently, in the amorphous phase only molecules of worse oriented kinetic segments remain. The mechanism of development of the amorphous orientation largely consists of uncoiling and straightening molecules under the effect of internal tensile stress in the fiber that accompanies stretching. The attempts at a theoretical description of the amorphous orientation development process consisting inits approximationby the affine deformation of the polymer lattice stretched unidirectionally presented by Kuhn and Grun [29] and Roe and Kringbaum[30,31] do not yieldresults adequate to the effects obtained for stretching PET fibers. The main cause is that PET fiber stretching is done with pseudo-affine deformation, not with affhe deformation as suggested by Ward [32]. The amorphous orientation is considered a very important parameter of the microstructure of the fiber. It has a quantitative and qualitative effect on the fiber deformability when mechanical forces are involved. It significantly influences the fatigue strength and sorptive properties (water, dyes), as well as transport phenomena inside the fiber (migration of electric charge carriers, diffusion of liquid). The importance of the amorphous phase makes its quantification essential. Indirect and direct methods currently are used for the quantitative assessment of the amorphous phase. Intermediate methods include the earliest procedure based on Stein’s equation [33] and one based on Samuels’ equation [34]. Among the direct methods is an IR spectroscopic method based onthe measurement of the dichroic ratio (R), of “amorphous” absorption bands. In the investigations [35], the amorphous bands 898 cm” and 1368 cm”, for which the angles of transition moment are a 8 9 8 = 39’ and a1368 = SO’, respectively, were used. Other methods are spectroscopy of polarized fluorescent radiation [35,36], measurement of color di-

Absorption band 875 cm- l 0.176 0.566 0.418 0.626 0.571 0.703 0.585 0.875 0.727

Absorption band 1343 cm- *

Absorption band 1473 cm-

0.269 0.615 0.691 0.817 0.870 0.921

0.309

0.851

chroismus of the fiber dyed according the test methods [37-391, the x-ray method, which consists of an analysis of the azimuthal distribution of intensity of the x-ray scattered diffusively on the fiber [40,41]. The values of the amorphous orientation index in deterthe form of Hermans’ function of orientation mined by the authors for PET fibers are listed in Table 6. The differences in the values offu quoted in Table 6 and referringto particular investigation methodscan result from the fact that in some methods the orientation of the amorphous and mesophase are consideredjointly. Consequently, in such a case the values of fa will be overrated.

vu),

3. OverallOrientation Overall orientation is understood as the joint arrangement of all the structural elements of the crystalline phase and noncrystallinepart of the fiber in relation to the geometrical axis of the fiber. In its essence, the overall orientation of PET fibers, as a result of the crystalline andamorphous orientation, will be characterized by smaller values of the quantitative index of orientation than for the crystalline phase and by greater ones for the amorphous phase. The quantitative assessment of the overall orientation of PET fibers is generally made on the basis of fiber optical anisotropy measurements, i.e., measurements of the optical birefringenceof the fiber. The determination of the value of optical birefringence makes it possible to determine the value of Hermans’ function of orientation based on the equation:

where Anideal is the optical birefringence of an ideal fiber, i.e., the fiber of a crystalline density and of ideal orientation, d, is the density of the crystalline region, A n is the birefringence, andd is the density of the fiber under investigations. The overall orientation of PET fibers depends on the production conditions and on further refining treatment.

Jeziorny848

and

Urbahczyk

Table 6 Amorphous Orientation Function (fa) of Differently Drawn PET Fibers

IR spectroscopy method

Draw ratio

x-ray method

2.0 x 3.0 x 3.5 x 4.0 x 4.5 x 5.2 x

0.208 0.645 0.350 0.744 0.816 0.416 0.834 0.510 0.886 0.599

Absorption band 898 cm- I 0.220 0.274 0.379 0.683 0.703 0.730

Absorption band 1368 cm"

Optical method a

0.496 0.362 0.522 0.784 0.642 0.834 0.681 0.898 0.780

b

0.605 0.715

Fluorescent polarization method 0.082 0.282 0.336 0.350 0.323 0.489

DRSDF methodc 0.047 0.089 0.156 0.162 0.130

-

An, = 0.253, An, = 0.220 1441. An. = 0.216, An, = 0.210 [45]. c Dichroic ratio of standard dyed fiber.

a

The factor having the strongest effect is the elongation imparted in the process of production stretching. Second, the overall orientation is affectedby the stretching rate. For the same draw ratio, the overall orientation grows withan increase in the stretching rate. The effect of the draw ratio on the value of Hermans' function of orientation is illustrated bythe values offo, established by the authors and depicted in Table 7. 111.

PHYSICALPROPERTIES

A.

MechanicalProperties

Among the basic mechanical properties of fibers are their deformability and tenacity. When an axial stretching force is appliedto the fiber, the principal quantitative indices of deformabilityare the axial elastic modulus (E)

Table 7 Overall Orientation Function (fo) of Differently

Drawn PET Fibers Draw ratio 0

2.0 x 3.0 x 3.5 x 4.0 x 4.5 x 5.2 x a

Hermans method a

b

Cunningham methodc

0.0164 0.1581 0.6709 0.7256 0.8169 0.8501 0.9057

0.0172 0.1657 0.7028 0.7601 0.8558 0.8906 0.9488

0.0149 0.1422 0.5924 0.6264 0.7140 0.7417 0.7918

Hermans method by assuming Anidea = 0.220 after Dumbleton

1441. c

Hermans method by assuming A nidea = 0.210 after Okajima 1451. Cunningham method [60].

(GPa) and the relative elongation at break ( E , ) (%). The first of the indices characterizes the resistance of the fiber to deformation under the affect of the acting force and determines the value of the tensile strength causing the formation of the unitary relative elongation ( E ) . The modulus E characterizes the susceptibility of the fiber to deform based on the principle of inverse proportionality. The other parameter, i.e., E,, characterizes the maximum ability of the fiber to deform under the effect of the acting tensile force. As for the tenacity, which describes the fiber's resistance to the action of the tensile force, the quantitative index assumed most often is the tensile strength (as) (cNltex). For the quantitative assessment of tenacity an index of mechanical long life ( 7 ) is used, proposed by Zurkow and Abasow[46,47]. This index definesthe time after which, when applying a definite tensile strength and a definite temperature, the fiber breaks.

1 . Axial Elastic Modulus ( E ) The axial elastic modulus of PET fibers ( E )depends, as with other kinds of fibers, on the value of the elastic modulus of crystalline material(E,) and amorphous fiber material ( E J . When approximating their fine structure by the phenomenological modelof Takayanagi,justified for PET fibers, it can be assumedthat the elastic modulus ( E ) for PET fibers is described by the equation for the Takayanagi model [48]:

(11)

where E, is the elastic modulus of the crystalline region. This is a constant value, characteristic for a given kind of space lattice of crystalline regions. For PET fibers, E, = 140 GPa according to Pierepelkin [49], E, = 110


849

GPa according to Sakurada and Kaji [50], and E, = 108 GPA according to Kunugi et al. [51]. The elastic modulus of amorphous material is E,, A and @ are parameters correspondingly proportional to the fraction of crystalline phase in the fiber. The product ( A . @ ) ,is defining the volume degree of fiber crystallinity (X); X = A.@. As stated by one of the authors [52], the parameters A and @ can be defined by the following expressions:

+ = -l

L A = -L X X l

b=E,X

c=EXE,X

a x E + b x E , + c = O

(14)

where: a = -l - X L '

l -EX-, L

1"

The value of E, in this case is a positive value of the root of the above equation. In the case of the lamellar substructure, i.e., for A = 1, typical of stretched and then annealed fibers, the equation has the form:

E, =

where l is the length of the crystallite; L is the mean long period; and X is the volume degree of fiber crystallinity. The parameter A determined by the above relationship can be calleda substructure parameter. Its numerical value determines the type of substructure of PET fibers. For A < 1 we have to do with the substructure of microfibrillartype, while for A = 1 with the substructure of lamellar type. Contrary to widespread opinion, the value of E, is not a constant quantity. As was proved previously[52], the value of E, is variable, since it depends on the ordering of macromolecules in the amorphous material of the fiber. Atthe same time, one can suppose that this ordering willbe affectedby the specificity of the fine structure of the fiber, and particularly bythe type of substructure of the fiber. The relationship determining the modulus E, appropriate for a definite type of fiber substructure can be derived from Eq. (1 1) when appropriate values of A are assumed. In the case of the microfibrillar substructure, i.e., forA < 1, typical of PETfibers stretched, but not subjected to annealing, this equation has the form [52]:

L1 [ [ L1 X+1"

EXE,X(l-X) E,-XXE

(15)

In the studies carried out by one of the authors [52], the values of E, and E were determined for PET fibers of the microfibrillar and of the lamellar substructure. The results have been presented in Tables 8 and 9. The results obtained showthat for both types of substructure the resistance to deformation, that is, the value of E, depends on the degree of molecular orientation of the amorphous material of the fiber V,) and the density of this amorphous phase of the fiber (d,). However, this dependence assumes a different formfor the microfibrillar and for the lamellar substructure. In the first case, it has the form: . = A X E 7

(16)

In the case of the lamellar substructure: da

E=CX-+D

(17)

fa

where A, B, C, D are corresponding constant values.

2. Elongation ut Break (E,) The elongation at break of PET fibers depends decisively onthe quantitative fraction and on the manner of the arrangement of the amorphous material of the fiber. In light of the results of investigations [53] that are presented in Table 10, one can assume that the value of the

Table 8 Fine Structure Parameters and Axial ElasticModuli of PET Fibers of Microfibrillar Substructure

Density of modulus Long Amorphous amorphous the orientation Degree material Crystallite period Draw (da) length function V a1 (1,) (nm) (&m3) ratio 3.0 x 3.5 x 4.0 x 4.5 x 5.2 x a

1.333 1.337 1.338 1.337 1.338

Determined by dsc method.

0.645 0.744 0.816 0.834 0.886

7.9 8.0 7.8 7.8 7.9

elastic of phase ( L ) Fiber parameter crystallinity (nm) (Xa) 11.2 12.0 13.4 14.0 16.6

0.36 0.36 0.37 0.38 0.39

Axial Substructure

Amorphous

(A)

(E) (GPa)

( E a ) (GPa)

0.51 0.54 0.64 0.68 0.82

9.1 9.2 9.7 9.9 10.8

4.4 4.7 5.5 5.6 6.5

850

Urbanczyk and Jeziorny

Table 9 Fine Structure Parameters and Axial Elastic Moduli of PET Fibers of Lamellar Substructure

Annealing temperature

ma

Density of the amorphous Amorphous Long material orientation Crystallite period (dU) function length (L) (g/cm3) (fa) (L)(nm) (nm)

120 140 180 220 a

1.337 1.330 1.333 1.345

0.310 0.265 0.185 0.115

4.6 5.1 5.9 7.6

Axial elastic modulus Degree of crystallinity (Xb)

Substructure parameter

0.44 0.46 0.50 0.56

10.4 11.0 11.8 13.8

(A 1

Fiber (El (GPa)

Amorphous phase (El)(GPa)

0.99 0.99 1 .oo 1 .oo

9.1 9.9 11.0 14.0

5.3 5.6 5.8 6.6

The fibers of draw ratio 1.5 x were annealed unstressed in the dryer for 2 h in air Determined by dsc method.

elongation at break ( E r ) increases with an increase in the fraction of the amorphous phase and a decrease in the degree of its orientation. It seems that at the same time both parameters act in the cumulated manner. Studies have shown the revealed straight line dependence between the elongation (€,) and the value of the cumulated index: [(l - X)/f,I4, where 1 - X is the fraction of the noncrystalline phase, while f a is the amorphous orientation function. In light of the above equations, the empirical dependence between the elongation at break and the fraction amorphous phase and its degree of orientation assumes the form:

vc),

tion function and the amorphous orientation function f a , Table 10. At the same time it appears that the enumerated parameters of the fiber fine structure, in the case of PET fibers, determine their tensile strength in the cumulative manner. The straight line dependence of the tensile strength on the product of p, fc, and fa, determined for PET fibers can testify to this fact [53]. The straight line correlation dependence being disclosed leads to the following empirical dependence of the tensile strength as a function of p, fc, and fa: u s

=

H x (p x

fc

x

fa)

+K

(19)

where H and K are constant values.

4 . Mechanical Long Life (7) According to Zurkow and Abasow [46,47], authors of this r index, it is determined by the relationship:

where F and G are constant values.

3. Tensile Strength r = TO

The tensile strength of PET fibers depends on their superstructure and internal orientation. The results from the investigations by one of the authors [54], show that the value of the tensile strength (as) is affected by the fraction of taut tie molecules (p),the crystallites orienta-

exp

[

u o - y x a R X T

1

where: TO = a constant characterizing the thermal vibrations of atoms inside the molecules;

Table 10 Structural Determinants of Elongation, Tensile Strength, and Mechanical Long Life of PET Fibers

Draw ratio

Fraction of noncrystalline material (1 - X)

Amorphous orientation function

Crystallites orientation function

Fraction of tie molecules

Elongation at break

(fa)

VC)

(P)

( E ~ )(%)

(cN/tex)

Structural parameter of the index T ( y ) (kJ/ molMPa)

3.0 x 3.5 x 4.0 x 4.5 x 5.2 x

0.64 0.64 0.63 0.62 0.61

0.645 0.744 0.816 0.834 0.886

0.817 0.820 0.83 1 0.842 0.844

0.027 0.030 0.040 0.043 0.055

52.9 26.7 17.3 16.8 14.5

19.2 30.8 42.7 55.3 77.0

0.380 0.242 0.174 0.134 0.096

Tensile strength (us)

Mechanical long life

(9

(7)

0.94 0.98 1.05 1.10 1.19


851

Uo (kJ/mol) = the activation energy of the breakdown of the fiber corresponding to chemical bond energy for PET fibers Uo = 175 (kJ/mol); Ro = gaseous constant = 8.314 x 10-3(kJ/mol x O K ) ; T ( O K ) = the fiber temperature; y (kJ/mol x MPa) = the structural parameter of the index r , and (T = applied tensile strength. The results of determinations of the T index for PET fibers obtained [54], have been listed in Table 10. In light of these results, one can state that the mechanical long life of PET fibers is determined by the specificity of their amorphous regions. The degree of molecular orientation and the density of packing of macromolecules in the amorphous material of the fiber must be considered the most important parameters. The mechanical long life (7) increases with an increase in the degree of molecular orientation of the amorphous phase. The increase in r is also promoted by looser packing of the macromolecules in this part of the fiber; i.e., its smaller density. These dependences are reflected in the empirical straight line relationship described by the equation:

where M and N are constant values. B. Selected Thermal Properties

2 . Thermal Volume Expansivity PET fibers, among basic synthetic fibers, distinguish themselves by a relatively low thermal volume expansivity. The value of the coefficient of the thermal volume expansivity ( a ) is dependent on the fiber temperature. For temperatures below the glassy temperature, the values of a are lower and, depending on the draw ratio of K-l, the fiber, are within the range 7.9-15.9 x while above the glassy temperature in the range K-'. The thermal expansivity anisot11.8-21.1 x ropy occurring in the fibers causes the a to be the resultant of the linear expansivity a11and aI , which quantita-

tively characterize the expansivity in the direction of the fiber axis and perpendicular to it: a = a11

+2x

a!I

(22) In the case of PET fibers, a11assumes the values close to 0, often even negative ones, aI ,which is perpendicular, is several times greater and always positive. AccordK-', and ing to Choy, et al. [SS], all = 0.5-0.3 x ayI = 4.2-7.8 x K - ' for the temperatures T < T,, andall = -0.5-0.5 x K-', anda. = 6.2-10.3 x K - ' for the temperatures T > T,. In the investigations carried out by one of the authors [56], an attempt was made to examine the conditions for the thermal volume expansivity of PET fibers. Within the framework of these investigations, a 2 5 was determined from the hydrostatic weight measurements using n-heptan as a liquid. The sought a 2 5 values have been calculated from the equation:

where Qo and Qt refer to fiber density at to = 19°C and t = 25"C, respectively. Fibers of a diversified draw ratio in the range 2.0-5.2 x were considered, determining the following parameters of their fine structure: the crystalline and amorphous orientation functions, fc and fa, degree of crystallinity, X , and critical dissolution time (CDS) in seconds. The results obtained are listed in Table 11. The a values obtained show that the thermal volume expansivity of PET fibers changes only slightly as a function of their stretching degree. At the same time, it can be observed that at the initial stage of stretching, i.e., for the draw ratio up to 3.5 x , the value of a deceases. For stretching above 3.5 x , they are higher and practically constant. It can be assumed that the variation of a is directly related to the changes in the coefficients of the linear expansivity, all and a. . As was stated by Choy et al. [55], at the initial stage of drawing of PET fibers, i.e., for the draw ratio below 3.5 x , there occurs a rapid decrease in allwhile simultaneously there is a very slow increase in Q.. Thus, the resultant outcome must lead to a decrease in the value of a. In turn, for advanced

Table 11 Thermal Volume Expansivity ( a ~and ) Related Fine Structure Parameters of Differently Drawn PET Fibers

Draw ratio 3.0 x 3.5 x 4.0 x 4.5 x 5.2 x a

a x 10-4 x 3.2 2.9 3.1 3.0 3 .O

K-1

fc

fa

X

TTM fraction

CDT" (s)

0.8 17 0.820 0.83 1 0.842 0.844

0.496 0.522 0.642 0.681 0.760

0.22 0.27 0.31 0.33 0.35

0.027 0.030 0.040 0.043 0.055

28 76 214 138 162

CDT = critical dissolution time in mixture phenoltetrachlorethan 2: 1 at 60°C.


852

drawing, i.e., for the draw ratio above 3.5 X, the rates of decrease in all and increase in a.,.compensate each other, which will lead to the stabilization of the value of a. It can be noted that the stated changes in a do not correlate with the variation of any separate considered fine structure parameter. This leads to the conclusion that the ascertained alteration in a must be evoked by a complex change in different parameters of the fiber substructure, microstructure, and orientation. The thermal volume expansivity of PET fibers undergoes an essential alteration resulting from annealing. The influence of annealing of PET fibers performed in air in a loose state at different temperatures and differentiated durations of heat treatment was studied [56]. The results obtained are presented in Fig. 7. As can seen, be the change in a depends on the annealing temperature and the duration of the heat treatment. The character of the alterations isattributed first to the level of the annealing temperature. Annealingtimemodifiesonly the range of changes butdoes not alter the character of the variation. The results obtained give evidence that there may be three distinguished temperature intervals of annealing in which a values attain different levels and undergo different alternations with an increasing annealing temperature. The first interval refers to the temperature range of13O-150/16O0C within which the a values are larger than those of unannealed fibers and monotonically decline with the growth of the annealing temperature. Thesecond temperature interval refers to 170-190°C

L f

20

I

110

I

160

I

180

Figure 7 Thermal volume expansivity coefficient of annealed PET fibers.

I

200

I

oc

within whichthe a values are lower than those of unana mininealed fibers and continuously diminish reaching mum value at 190°C. The third temperature interval pertains to the temperature range of 200-210°C. Within this zone a increases with the annealing temperature reaching a level equal to or even higher than that of unannealed fibers. Comparingthe alteration in a values with the established changesin fine structure parameters presented in Table 12, it can be seen that the alteration in a values does not correlate with variations in any particular structure parameter. This leads to the conclusion that the ascertained alternation in a must be evoked by a complex change in different substructure, microstructure, and orientation parameters of the fiber. 2. Heat Capacity

PETfibers distinguishthemselves among basic synthetic fibers by a relatively small value of heat capacity. The value of the heat capacity coefficient at constant pressure lies in the range of C, = 0.24-0.44 (cal K" g"). It seems that the relatively low C, value of PET fibers reflects the facility in enhancing the fibers internal energy by conveying heat to them. The heat capacity of PET fibers undergoes essential alternation as a result of the applied heattreatment of the fiber. The relationship between the c, and the annealing conditions of PET fibers has been studied [561. The heat capacities of a 4.5 x drawn filament were estimated at constant pressure from DSC diagrams using the relationship:

where H, and H, refer to the amplitude of enthalpy change occurringat the temperature increase from 20 to 25"C, respectively, for fiber sample and reference substance (saphire); m, and ma are weights of sample and reference substance; and c, = 775.43 (JkGK) is the heat capacity of the reference substance. The obtained c, valuesand the established fine structure parameters for the PET fibers examined are presented in Fig. 8 and Table 13. The results obtained show a certain general regularity. The c, values in the range of an annealing temperature of 130-160°C are higher than for the unannealed fiber and increase with an increase in temperature. For the range of temperatures exceeding 160°C,the c, values monotonically decrease and assume smaller valuesthan for the unannealed fibre. Comparing the alternation in c, with changes in the established fine structure parameters, it will be noted that the change in c, does not correlate with the variation of any separate structure parameter. Closer examination of the relations enables us to infer that the c, values of annealed PET fibers are determined by a counteracting, mutually competitive influence of particular structure parameters. The level

Fibers

Fine

of PET

853

Table 12 Fine Structure Parameters of Annealed PET Fibers

Annealing crystallinity Birefringence Annealing temperature

“M (min) time

(“C)

Without 130

140

160

180

200

210

-

2 5 10 30 2 5 10 30 2 5 10 30 2 5 10 30 2 5 10 30 2 5 10 30

(An )

An Id

0.1286 0.1681 0.1615 0.1614 0.1653 0.1645 0.1607 0.1593 0.1670 0.1628 0.1572 0.1628 0.1622 0.1592 0.1610 0.1606 0.1650 0.1637 0.1601 0.1634 0.1663 0.1623 0.1611 0.1573 0.1692

0.1296 0.1209 0.1165 0.1184 0.1189 0.1184 0.1156 0.1146 0.1202 0.1169 0.1129 0.1170 0.1163 0.1139 0.1154 0.1150 0.1180 0.1171 0.1145 0.1167 0.1187 0.1158 0.1150 0.1122 0.1203

(xa)

( W

36.7 45.6 45.8 45.4 49.2 48.1 48.6 49.0 48.8 50.3 50.8 51.3 52.8 53.5 53.1 54.3 55.6

55.5

55.4 57.8 57.9 58.7 58.1 59.5 62.3

dissolve

time fraction

(S)

orientation function (L)

9

0.107 0.048 0.047 0.045 0.036 0.045 0.041 0.043 0.036 0.043 0.041 0.043 0.034 0.041 0.038 0.040 0.029 0.038 0.038 0.039 0.029 0.021 0.031 0.018 0.021

0.50 0.50 0.50 0.47 0.46 0.49 0.49 0.46 0.45 0.47 0.44 0.44

19 25 41 55 40 35 55 74 103 108 106 156 160 150 222 245 240 245 323 383 329 405 703 1417

0.44

0.46 0.44 0.44 0.43 0.45 0.44 0.44 0.43

0.44

0.43 0.43 0.42

of c, is always set by mutualopposite effects involved by the change of taut tie molecules fractionand the overall orientation on one hand,and the change in crystallinity on the other.

Table 13 Heat Capacity of Annealed PET Fibers

(min)

Time

Annealing temperature (“C)

0.1

t 20

160

160

180

200 OC

Figure 8 Heat capacity c,, of annealed PET fibers.

130 140 150 160 170 180 190 200 210

Without

2

30

0.412 0.421 0.405 0.398 0.332 0.323 0.305 0.151 0.140 0.269

5

10

0.363 0.453 0.442 0.428 0.415 0.241 0.192 0.167 0.163

0.3% 0.429 0.452 0.474 0.380 0.346 0.254 0.226 0.216

-

-

0.376 0.441 0.433 0.344 0.340 0.227 0.194 0.090 0.097

-


854

C. SelectedElectricalProperties

l . ElectricalConductivity The electrical conductivity of PET fibers as compared with other main synthetic fibers is relatively low. This explains why PET fibers are often utilized inthe manufacture of textiles as electroisolating materials. The value of the electrical resistivity characterizing recipro(Oecm). The mechcalconductivityis of the order anism ofthe electrical conductivity of PET fibers is still a matter of controversy. According to results attained [57], there are convincing arguments that in the case of PET objects the electrical conductivityis due to the ionic mechanism. The low electrical conductivity of PET fibers depends essentially on their chemical constituency, but also to the same extent on the fiber's fine structure. In one study [58], an attempt was made to elucidate the influence of some basic fine structure parameters on the electricalresistivity of PET fibers. The influence of crystallinity (x)the average lateral crystallite size ( A ) , the mean long period ( L ) ,and the overall orientation function WO)have been considered. The results obtained are presented in the form of plots in Figs. 9-12. The reported results give evidence that the electrical conductivity decreases with increasing crystallinity and with enhancement ofthe mean long periodof the fiber. In addition, the conductivity simultanously decreases with the diminishing of the overall orientation. It seems that this relationship can be explained onthe basis of the nature of electrical conductivity, i.e., on the migration of carriers of free electrical charges. The increase in crystallinity accompanied by the increase in L create more hindrances for the migrationof carriers of electrical charges and, therefore, impedes the conductivity. The opposite result is caused by the improvement of the overall orientation.

d 0,110

0,150

0,130

BIREFRINGENCE

h

.

Figure 10 Electrical resistivityof PET fiber versus birefringence of the fiber.

C nocm 3

1012 W

a

d

: c

U W

1011 1010

-1

W

16

24

32

CRYSTALLITE SIZE

L nm 3

D(100)

Figure 11 Electrical resistivityof PET fiber versus average crystallite size perpendicular to the chain direction.

>. c CI

C n*cm 3

2c I? cn W a

d a c

U

40

50

60

70

C X 3

DEGREE OF CRYSTALLINITY

X

Figure 9 Electrical resistivityof PET fiber versus degree

of crystallinity of the fiber.

i

11,o HEAN

12,o

13,O

C nn J

LONG SPACTNG L

Figure 12 Electrical resistivityof PET fiber versusmean long spacing.


855

2. ElectrostaticCharges PET fibers are characterized by a strongly pronounced liability to electrostatic charging. Their high susceptibility to electrostatic charging is determined by a polar character of the polymer, and their low electrical conductivity. Investigations carried out by the authors [58] indicate that the effect of electrostatic charging of PET fibers depends additionally on the fine structure of the fiber. In these investigations an assessment was made of a charge in the surface density of electrostatic charges ( F ) (C/cm2), generated during the rubbing of the fiber against a brass cylinder as a function of the crystallinity degree, average lateral crystallite size, mean long period, and overall orientation. The results obtained have been presented in the form of plots in Figs. 13-16. The dependences shown indicate that F increases with an increase in the crystallinity degree. The dependence on the average lateral crystallite size and the overall orientation is of the curvilinear character. Based on the results attained, it can be concluded that the electrostatic charging of PET fibers will be particularly strong for fibers of a high degree of crystallinity, coarse-grained crystalline structure, and low overall orientation. D. Optical Properties

L C/cn2

to-B 1o

0,110

0,150

0,lTO

BTREFRINGENCE An

Figure 14 Electrostatic charge of PET fiber versus birefringence of the fiber.

3 a

c

C/ctnL

3

a X

U

tl-!

1o

t-

The basic opticalproperties of PET fibers consist of the anisotropy of light transmission and the anisotropy of light absorption. The first one manifests itself by the light doublerefraction, the other by the pleochroic effect of the fiber coloring. The double refraction causes the refractive indices of the ordinary and extraordinary light beams (no,n,) originated by light transmission through

- ~

lo-1o

a

'

3

z

- ~

0,

t-

U W -l W

10-l0

16

24

S2

CRYSTALLITE SIZE

t nm 3

D(100)

Figure 15 Electrostatic chargeof PET fiber versus aver-

age crystallite size perpendicular to the chain direction.

c

c/cm2

3

c

10-

10'~

-10

10

40

I

I

I

SO

60

70

J

I

C X 3

DEGREE CF CRYSTALLINTTY

X

Figure 13 Electrostatic charge of PET fiber versus fiber degree of crystallinity.

11.0

12.0

13.0

PIEAN LDNG SPACING

C nm 3 L

Figure 16 Electrostatic chargeof PET fiber versusmean long spacing.

ratio

Urbahcryk and Jeriorny

856

Table 14 Directional Refractive Indices (rill, nl) and

Birefringence ( A n ) of Differently DrawnPET Fibers Draw 1.0 x 2.0 x 3.0 x 3.5 x 4.0 x 4.5 x 5.20.1881 x 1.5343

ne

nl

An

1S732 1S934 1.6866 1.7037 1.7119 1.7172 1.7224

1 S696 1 S648 1 S470 1 S444 l S409 1 S382

0.0036 0.0286 0.1396 0.1593 0.1710 0.1790

the fiberto differ in their values. In view the of fact that PET fibers are optically positive, ne is always greater than no. Considering that ne is the value ofthe refraction index for thewave vibrating parallelto thefiber axis, it is additionally denoted by the symbol rill. In turn, no being the value of the refractive index for thewave vibrating perpendicular to the fiber axis is also denoted by nl. The measure of the optical anisotropy of light transmission of the fiber is optical birefringence ( A n ) .

A n = ne - no

nil - nL

(25) In the case of PET fibers, the main directional refractive indices rill, nL and the optical birefringence A n are dependent on the magnitude of the draw ratio at which the fiber was obtained and to a lesser extent on its crystallinity degree. The values ofrill, nL ,and A n for fibers of a diversified draw ratioare shown in Table14. The ability to refract light independent the of direction of the light wave vibration is presented by the isomorphic refractive index (ni,,). For PET fibers niso = 1.598. In the case of PET fibers, we have to deal with the occurrence of internal radial differentiation of the birefringence. Results from investigations performed by one of the authors [59] show that the radial differentation of the A n values varies for fibers of different stretching. For nondrawn fibers, the following are characteristic: a low value of A n for the central axial part of the fiber and a low valuefor thesurface layer.The remaining part of the fiber is distinguished by higher and, at the same time, variable values ofA n . Starting withthe draw ratio 2.0 x , the value of A n for thesurface layer and the central part increases. For the highest draw ratio (4.5-5.2x),the higher values of A n occur only for the surface layer.

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6. D. C. Prevorsek and Y. D. Kwon, J . Macromol. Sci., 852: 447 (1976). 7. A. Jeziomy, Acta Polymerica, 41: 590 (1990). 8. R. T. Daubeny and C. W. Bunn, Proc. Roy. Soc.,A226: 531 (1954). 9. Ju. Ja. TomaszpolskijandG. S. Markova, Vysokomol. Sojed., 6: 274 (1964). 10. S. Fakuov, W. Fischer, and G. E. Schmidt, Makromol. Chem., 176: 459 (1975). 11. C. Yamashita, J . Polym. Sci., A3: 81 (1%5). 12. J. L. Koenig and M. L. Hannen, Macromol. Sci.-Phys., BI: 617 11971). . -,13. K. K .Mocherla and J. P. Bell,J . Polym. Sci.-Phys. Edn., 11: 1779 (1973). 537 14. S. A. Baranova et al., Vysokomol.Sojed.,A22: (1980). 15. S. L.Pakszwer et al., Vysokomol. Sojed., A25: 37 (1984). 16. T.Asano and T.Seto, Polymer J., 5: 72 (1973). 17. W. L. Lindner, Polymer, 14: 9 (1973). 18. G. Urbariczyk and I. Waszkiewicz, Polimery, 25 (253): 440 (1980). 19. A. Jeziorny, Zesz.Nauk.P.L.-Inz.Wlok., 2116: 195 (1993). 20. Z.W. Wilchinsky,Adv. X-Ray Anal., 6: 231 (1963). 21. J. H. Dumbleton and B. B. Bowles, J . Polym. Sci., A24: 951 (1966). 22. V. B. Gupta and S. Kumar, Text. Res. J., 49: 405 (1979). 23. G. Urbanczyk, J . Appl. Polym. Sci., 51: 201 (1994). 24. R. D. B. Fraser, J . Chem. Phys., 21: 1511 (1953). 25. R. D. B. Fraser, J . Chem. Phys., 28: 1113 (1958). 26. W. A. Chranowski, Vysokomol. Sojed., B25: 96 (1983). 27. G. Urbariczyk, Polimery, 21: 26, 69 (1976). 28. G. Urbanczyk,J . Polym. Sci.- Symp., 58:311 (1977). 29. W. Kuhn and F. Gnin, Koll. Z., 101: 848 (1942). 30. R. I. Roe and W. R.Kringbaum, J . Appl. Phys. 6: 32 (1962). 31. R. I. Roe and W. R. Kringbaum, J . Appl. Phys.35: 2215 (1964). 32. I. A. Ward, Structure and Properties of Oriented Polymers, Appl. Sci. Publ. Ltd., London, p. 34 (1975). 33. R. S. Stein and F. H. Noms, J . Polym. Sci., 21: 381 (1956). 34. R. I. Samuels, J . Polym. Sci., A2-3: 1741 (1%5). 35. G. Urbadczyk, Zesz. Nauk. P. L.,nr.646, Wlokien., 50: 83,55, (1993). 36. G. E. McGraw, J . Polym. Sci., A2-8: 1323 (1970). 37. J. M. Preston, J . Soc. Dyers Colour, 66: 361 (1950). 38. J. G. Blacker and D. Patterson,J . Soc. Dyers Colour, 85: 598 (1968). 39. B. Lipp-Symonowicz and G. Urbariczyk,-Polimery, 2 7 198 (1982). 40. H. J. Biangardi, J . Polym. Sci.-Phys., 18: 903 (1980). 41. A. Jeziorny, Polimery, 39: 234 (1994). 42. W. Ruland, Polymer, 5: 89 (1964). 43. G. Vonk, J . Appl. Crystallogr., 6: 148 (1973). 44. J. H.Dumbleton, J . Polym. Sci., A-216: 795 (1%8). 45. S. Okajima and K. Kayama, Sen.Z.Gakkaishi, 22: 51 (1976). 46. S. N.'Zurkow and S. A. Abasow, Vysokomol. Sojed., 3: 441 (1961). 47. S. N. Zurkow and S. A. Abasow, Vysokomol. Sojed., 4: 1703 (1962). 48. M.Takayanagi, K. Imada, and T. Kajima, J . Polym. Sci., CIS: (1966). 49. K. E..Pieriepielkin, Chim. Volok., 2: 3 (1966). 50. I. Sakurada and K. Kaii, J . Polvm. Sci.. C31: 57 11989). ~

.

.

Fine

of PET Fibers

51. T. Kunugi, A. Suzuki, and M. Hashimoto, J . Appl. (1988). Polyrn. 84 Sci., (1981). 26: 1951 52. A. Jeziorny, Przegl. Wlokien., 47(10): 235(1993). 53. A. Jeziorny, Przegl. Wlokien., 48(10): 3(1994). 54. A. Jeziorny, Przegl. Wlokien., 49(4): 3 (1995). 55. C. L. Choy, M. Ito, and R. S. Porter, J . Polyrn. Sci-Phys. (1983). Ed., 21: 1427 56. G. Urbaiiczyk and G. Michalak, J . Appl. Polyrn. Sci.,32: (1974). 4787, 3841 (1986).

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857 57. G . Urbaiiczykand W.Urbaniak-Domagala, Polirnery,33: 58. G . Urbaiiczyk and A. Jeziorny, Faserf. U. Texfil., 24: 151 (1973). 59. G. Urbdczyk* zeSZYo Nauk* p - 1.646, W1ok*50:

60. F. Cunninghamand G. R. Davies, Polymer, 15: 743


Index

activated swelling procedure, 222 activated systems, 42 activation energies, 735 activation energy, 127, 561, 779 activation of the neuron, 3 activation parameter, 265 active groups, 130 active sites, 760 acylation, 268 acylation reaction, 261, 267 addition reactions, 736, 740 additives, 140, 141, 142 adhesion, 272, 822 adhesion capability, 270, 275 adhesion properties, 270 adiabatic polymerization, 66 adsorption, 370 adsorption equilibrium, 128 adsorption of anionic polymers, 128 adsorption rate, 207 adsorption theory, 823 aggregation behavior, 770 aggregation process, 127 aging time, 572 agricultural applications, 114 alcohols, 140 aliphatic compounds, 174 aliphatic cyclic amines, 233 aliphatic diisocyanate, 139 aliphatic solvents, 175 alkali fiber glass, 85 alkali metal cations, 348 alkali metal hydroxides, 86 alkali metals, 86 alkali resistance, 423

ABA block copolymer, 244, 765 ABA copolymer, 244 ablative materials, 428 ABS, 317 absorbing monomer, 212 absorbing UV light, 115 absorption, 132 absorption bands, 87, 130, 300 absorption of light, 243 accelerators, 379 acetic acid, 331 acetic anhydride, 266 acetone, 215 acetone extraction, 173 acetonitrile, 210 acetoxy groups, 330 acetylene, 495 acid catalysis, 331 acid catalysts, 424 acid concentration, 530, 535 acid dissociation, 99 acid functionalized rubber, 441 acid group, 145 acrylamide, 62, 65, 69, 124 acrylamide in aqueous solution, 124 acrylamide monomer, 251 acrylamide polymerization, 120 acrylamide polymers, 61, 65 acrylic esters, 164 acrylic monomer, 227 acrylicacid-functionalized polypropylene, 591 acrylonitrile, 332 activated aromatic dihalides, 37 activated difluoride, 54 activated swelling process, 215 859

860

alkaline earth cations, 339 alkaline hydrolysis, 345 alkenylation reaction of polystyrene, 266 alkoxy derivatives, 50 alkoxy radicals, 79, 492 alkylaluminum compounds, 322 alkylation of polystyrene, 264 alkylation reaction, 261, 263, 264 alkylcyanoacrylates, 210 alloying cycloolefin copolymers, 164 alloys, 649 alpha-olefin comonomer, 278 alternating copolymer, 481 alternating stress conditions, 149 alumoxane-free commercial catalysts, 155 American Society for Testing and Material, 442 amide, 36 amide groups, 129, 133 amine groups, 260 aminium radical, 234 aminium radical cation, 232 amorphous orientation, 846 amorphous phase, 844 amorphous polymers, 395 amorphous structure, 54 anacardic acid, 421 anacardium occidental, 421 analytical chemistry, 780 anhydride groups, 266 anhydrous conditions, 42 aniline, 238 animal fibers, 813 animal skins, 411 anion exchange, 133 anionic chain polymerization-anion radical transfer, 744 anionic emulsifier, 221 anionic grafting, 535 anionic polymerization technique, 602 anionic PVC, 323 annealing, 395 annealing process, 391 antifoaming, 769 antioxidant, 113, 470 antioxidant action, 105, 115 antioxidant mechanisms, 105 antioxidants, 117, 399, 768 antishrink efficiency, 581 anti-synergistic effects, 117 apolar monomer, 215 apparent viscosity, 449 application for polypropylene, 160

Index

applications of PBAs, 663 aqueous brines injections, 71 aqueous emulsion, 221 aqueous solution polymerization, 65 aqueous solutions of polymers, 62 architectural differences, 2 aromatic amines, 115, 239 aromatic compounds, 174 aromatic difluorides, 36 aromatic diisocyanates, 139 aromatic polyesters, 88 aromatic polyimines, 50 aromatic solvents, 174 aromatic systems, 36 aromatic tertiary amine, 228 Arrhenius equation, 127, 366 Arrhenius law for frequency-temperature shift, 393 arsonium ylide, 377 artificial neural network, 1 atactic blocks, 159 Atomic Energy Authority, 133 atomic structure, 24 attenuated total reflectance spectroscopy, 517 Aureobasidiae strains, 94 autoacceleration, 375 autoinitiation reaction, 105 automobile panels and bumpers, 652 auto-oxidation/oxidative coupling polymerization, 420 autoxidation cycle, 111 autoxidation theory, 105 average molecular weight, 210, 223 average pore size, 221 axial elastic moduli of PET fibers, 849 axial elastic modulus, 848 azeotropic water mixtures, 67 azo compounds, 119 azo groups, 502 azo initiator, 735 azo-containing polymers, 736 backbone relaxation, 443 backbone scission, 493 back-propagation, 8, 22 back-propagation algorithms, 20 back-propagation training, 24 band fine splitting, 98 barrier properties, 580 BASF (Cheadle, UK), 765 batch-off system, 454, 456 batch processing, 717 Batschinski's equation, 13

861

Index benzene, 172, 173, 175, 357 benzophenone, 239 BF Goodrich Chemical Group, 143 Bicerano's process, 26 Afunctional initiator, 742 binding energy, 521 biochemical properties, 100 biochemistry, 95 biocompatibility, 102, 508 biological activity, 179, 183, 185 biological antioxidants, 115 biological immuno-stimulating activity, 185 biologically degradable composite materials, 806 biomass, 413 biosynthesis, 94 biotechnology, 189 bisphenols, 36, 37, 54 bitumen, 418 blend compositions, 150 blend morphology, 477, 624 blend morphology/property relationship, 656 blending, 138, 142, 150, 651 blending behavior, 428 blending of ionomers, 152 blending of polymers, 649, 667 blending system, 140 blending with TPU, 140 blends, 152,471,623 blends miscibility, 589 blends of NBR-ACM, 614 blends of NBR-CSPE (Hypalon), 612 blends of NBR-thiokol, 615 blends of XNBR-AU, 619 block and graft copolymers, 725, 731 block copolymer of polypeptides, 253 block copolymer sequences, 654 block copolymer synthesis, 735 block copolymer synthesis with low-molecular weight azo compounds, 748 block copolymerization-systems, 759 block copolymers, 481, 602, 605, 634, 643, 668, 726, 735, 738, 746, 755, 760 block copolymers derived from macroinitiators, block copolymers of acrylamide, 69 block efficiency, 760 blow moulding, 278 blowability, 283 blown film, 139 BMI cured blend, 472 bond formation, 360 bond indices, 26

bonding agent, 571 bonding groups, 269 bone cement, 227 boric acid esters, 88 boric acid polyesters, 88 Bragg spacing, 145 brake linings, 423 brake linings formulations, 424 braking action, 424 branching, 73, 88, 320 break strain, 392 bromine, 429 bromobenzene, 175 Brookfield viscometer, 169 bulk polymerization, 232 bulk polymers, 318 bulk process, 656 bulk specimens of ionomers, 147 bulky cations, 739 calender, 142 calorimetry measurements, 54 carbanion, 373 carbohydrates, 413 carbon black, 399 carbon disulfide, 177 carbon-carbon bonds, 82 carbon-to-carbon double bonds, 31 carbonyl, 81 carbonyl index, 403 carboxamides monomers, 42 carboxyl group, 262, 268, 269 carboxylate terminus, 99 carboxylation reaction, 263 cardanol, 424, 427, 432 cardanol grafted cellulose, 431 cardanol possesses, 429 cast thin films, 150 catalysts, 155, 566, 776 catalyzed hydrolysis of esters, 776 cation carriers, 339 cation exchange resins, 775 cation radical transfer, 741 cationic catalysis, 259 cationic catalyst media, 267 cationic metallocene catalysts, 164 cationic poly acrylamide, 120 cationic polymerization, 427 cationic resins, 775, 777, 780 cationic resins from Bagasse Pith, 783 cationic species, 130 cationic techniques, 748

862

cationomer, 151 cations, 339 cell membrane, 180 cellular targets, 103 cellulose, 69, 173, 413, 416, 426, 535, 791 cellulosic filled plastic systems, 577 cellulosic materials, 583 cellulosics, 577 centrifugation, 288 eerie ion, 546 eerie ion initiation, 503 eerie ion redox, 541 eerie ion redox initiation systems, 541, 542 eerie ions, 530 chain branching, 278, 283, 284 chain branching in HP LDPE, 278 chain branching in LLDPE, 278 chain breaking antioxidants, 117 chain breaking donor antioxidants, 112, 113 chain breaks, 353 chain flexibility, 35 chain initiation, 125 chain length, 353, 356, 363, 426 chain polymerization, 750 chain radical, 529 chain reaction, 81 chain scission, 332 chain transfer agents, 379 chain transfer initiation, 502 chain-end unsaturation, 319 chain-folded crystallization, 296 chain-solvent interactions, 363 chalking, 110 change of surface tension, 796 characterization of interface, 826 characterization of PBAs, 655 charge transfer initiation systems, 227 charge-transfer, 236 charge-transfer polymerization, 227 chemical activation, 483, 536 chemical bond, 475 chemical bonding, 468 chemical compositions, 791 chemical coupling, 7% chemical degradation, 351 chemical derivatization, 99 chemical features of materials, 689 chemical methods, 502 chemical modification conditions, 275 chemical modifications, 145, 263, 415 chemical nature, 161 chemical properties, 63, 99, 815

Index

chemical reactions, 141, 669 chemical stability, 819 chemical stability of low-molecular compounds, 366 chemical thermodynamics, 363 chemical treatment, 823 chemically modified natural monomers, 418 chemorheology, 712 chlorinated PE, 640 chlorinated surface of glass, 255 chlorination of PVC, 325 chlorobenzene, 176, 177 chloroform, 347 chloromethyl group, 212, 218 chloromethylstyrene, 218 chlorometylation reaction, 260 chlorosulfonated polyethylene, 466 chromatographic analysis, 232 chromatography, 94 chromic acid etching, 520 chromism in organic solvents, 173 chromophore quenchers, 115 citric acid esters, 140 class of catalysts, 155 classification of composites, 811 classified PBAs, 653 cleavage methods, 767 clinical applications, 772 cloud point, 769 clutch facings, 423 (co)polymerization, 68 coal, 417 cocrosslinking, 465 coefficient of variation, 189 cohesion, 145 cold feed extruders, 457 cold feed extrusion, 456 cold pot, 462 cold press molding, 818 colloid coatings, 167 color change, 177 color development, 174 color of porous film, 175 combustion chamber, 705 commercial polymers, 35, 287, 322 commercial products, 142 commercial PVC resin, 138 commercial PVC/TPU polymeric blends, 143 comonomer, 216, 217 compatibility, 138, 139, 644 compatibilization, 633, 637 compatibilization concepts, 634, 667 compatibilization of NBR-PP, 647

Index

compatibilization of PBAs, 654 compatibilization process, 670 compatibilized composite, 592 compatibilizers, 635, 761 compatibilizing action, 594 compatibilizing an immiscible system, 591 compatilization of TPEs, 638 compleximetric methods, 768 components of CNSL, 423 composite properties, 819 composite sheet laminates, 579 composites, 650 composition of PVC/TPU blends, 138 composition property relationship, 611 compound mixing, 454, 455 compression molding, 52, 142, 459 computer prompted weighing system, 456 concentrated particle processing, 707 concentrated solutions, 17 condensation of ethylene oxide, 766 conformational disorder, 300 conjugated double bond sequences, 326 conjugated polyene structures, 320 connectivity indices, 26 consistency, 451 contact angle measurement, 518 continuous processing, 718 control of block sequences, 759 conventional plasticizer, 152 conventional radical initiators, 483 conventional swelling process,, 215 conventionally vulcanized elastomer-plastic blends, 469 cooling process, 311 cooling rate, 297 coordinate bonds, 132 copolyesters, 3% copolymer latex particles, 218 copolymer latices, 217, 218 copolymer macromolecules, 481, 482 copolymenzation, 69, 120, 145, 321, 329, 547, 732 copolymenzation of acrylamide, 69 copolymenzation of styrene, 218 copolymenzation reaction, 217 copolymerizing antioxidant functions, 116 copolymers, 127, 637 copolymers of acrylamide, 119 copolymers with higher alpha olefins, 160 core-shell type microspheres, 609 core-shell type polymer microspheres, 601 corona discharge treatment, 824 correlation angles, 386

863 cost effectiveness, 117 cotton, 411 counterion, 149 coupling agents, 799, 820, 825 covalent bonds, 146 covulcanization, 465 crack surface, 171 cracks, 310

crammer feeder, 721 craze fibral breakdown, 146 craze fibrillation, 151 craze initiation, 480 crazing, 146 critical chain length, 208 critical ion content, 145 crosslink density, 30, 31, 602 crosslink sites, 30 cross-linkable phosphorus containing polyimides, 46 crosslinkages, 365 crosslinked cellulose, 416 crosslinked graft copolymers, 733 crosslinked polymers, 420 crosslinked seed network, 215 crosslinked structures, 131 crosslinker, 220 crosslinking, 640, 825 cross-linking agents, 466 crosslinking reaction, 219 cryogenic method, 655 crystal structures, 295 crystalline dynamics, 295 crystalline homopolymer, 151 crystalline lamella, 840 crystalline orientation, 845 crystalline phase, 839, 841 crystalline polyethylene, 295, 331 crystalline polymers, 287 crystallinic fibrils, 793 crystal Unities, 161 crystallinity, 54, 148, 163, 284, 287, 325, 390 crystallites, 842 crystallites in microfibrils, 840 crystallizability, 102 crystallization, 2%, 297, 303, 669 crystallization curve, 297, 298 crystallization of FCC, 304 crystallization of PEs, 284 crystallization of polyethylene, 295 crystallization pressure, 304, 314 crystallization processes, 312, 395 crystallization temperature, 308, 314

864 cure enhancer, 428 cure rate, 466

cured polyethylene, 90 curing agent, 75, 379 curing systems, 470 cyclic compositions, 361, 369 cyclic imide structures, 64 cyclic transition state, 238 cyclobutene, 164 cyclohexane, 172, 174, 178 cyclo-olefin polymers, 164 cyclopentene, 164 cytokine induction, 187 cytosolic contents, 96 cytotoxic action on cultured cells, 184 cytotoxicity, 185 Darcy's law, 817 data processing, 8 DBM groups, 523 decafluorobiphenyl, 36 decomposition, 488, 492 deformable layers, 795 deformation modes, 146 degradation, 100, 318, 399 degradation effects, 30 degradation rates, 324 degradative agents, 317 degree of acylation, 267 degree of conversion, 148 degree of hydrolysis, 167 degree of neutralization, 148 degree of polymerization, 1% dehydrochlorination, 324 dehydrochlorination rate, 325 depolymerization, 82 destruction of polyolefins, 83 detergency, 770 dextran, 211 diacyl peroxide amine, 231 diacyl peroxide-amine systems, 227 dialkyl sebacate esters, 14 diallylammonium salts, 126 Dianippon Ink & Chemical Company, 143 dielectric relaxation, 512 dielectric thermal analysis (DETA), 393 diene rubbers, 448 differential scanning calorimetry, 138 diffusion, 192 diffusion rate, 199 diffusion theory, 823 diffusion-controlled reactions, 126

Index

difunctional unit, 10 dilatancy behavior, 169 dimensional stability, 651 dimethyl methylene-bridged cyclopentadienyl fluroneyl zirconium dichloride, 163 dimethyl siloxane rubber (PDMS)-based blends, 676 dimethylformamide (DMF) treatment, 141 dioctylphthalate (DOP), 137 diols, 738 dipolar compounds, 373 direct method, 508 direct reactive blend processing, 659 direct titration, 780 discomposition reactions, 89 disentanglement, 281 dispersed elastomenc micro-particles, 479 dispersed phase particle size, 658 dispersion, 476, 599, 770 dispersion copolymerization, 215 dispersion medium, 205, 207 dispersion polymerization, 67, 201, 202, 208, 210, 211, 220, 222, 223 dispersion polymerization of acrolein, 211 dispersion polymerization of styrene, 218 dispersion polymerization process, 210 dispersion polymerization recipes, 202 displacement polymerizations, 35, 39, 42, 50 displacement reaction, 35 disproportionate, 89 dissolved oxygen, 193 divalent cations, 340 divinylbenzene, 222 DMAC, 36, 46 DNA, 100, 102 DNA polymerase, 100 DOP, 140 dose rate effects, 509 double bonds, 87, 528 double irradiation method, 493 draw ratio, 597 drilling-fluids, 71 droplet deformation, 586, 589 dry adhesive bonding process, 579 dry bonding without adhesives, 579 DSC curves, 314 DSC melting curves, 312, 388 DTA curves, 297 dynamic behavior, 571 dynamic mechanical properties, 392 dynamic mechanical studies, 444, 447 dynamic swelling method, 215

Index

dynamic vulcanization, 467, 474 dynamic vulcanization techniques, 476 dynamically vulcanized blends, 474 E-P copolymer, 403 ebonite method, 655 ecology, 773 effect of grinding on the graft ability, 536 effect of pH, 129 effect of process variables, 203 effect of solvent, 779 efficiency of the copolymers, 635 elastic energy, 617, 618 elastomer, 27, 465 elastomer-elastomer blends, 466 elastomer toughened polycarbonate blends, 659 elastomer toughened styrene system, 656 elastomeric alloys, 653 elastomeric blends, 468 elastomeric compound, 451, 455, 459, 463 elastomeric sealants, 27 electrical conductivity, 854 electrical properties, 854 electro-neutrality, 191 electron acceptors, 236, 237 electron beam treatment, 524 electron donors, 236 electron microscope, 304 electron-attracting monomers, 227 electron-transfer, 86 electron-transfer processes, 236 electronic microscopy, 138 electrostatic theory, 823

elongation, 76, 110, 595, 651 elongational rate, 283 emulsification, 770 emulsifier, 68, 190, 196, 771 emulsifier concentrations, 197 emulsifier molecules, 190 emulsion copolymerization, 198 emulsion copolymerization reaction, 216 emulsion medium, 190 emulsion polymerization, 68, 190, 192, 193, 194, 195, 197, 198, 220 emulsion polymerization of styrene, 198 emulsion polymerization of VAc, 170 emulsion polymerization process, 191 emulsion polymerizations of styrene and vinyl acetate, 197 end group analysis, 232, 233 end use characteristics, 157 end use performance, 290

865

energetic composites, 705, 722 energy content, 707 energy transfer, 400 energy-to-fracture, 147 engineering thermoplastics, 56, 623 entanglement molecular length, 641 entanglement strand density, 152 entropies, 300 entropy of mixing, 138 environmental conditions, 94 environmental effects, 808 environmental stress cracking, 56 enzymatic hydrolysis by polymalatase, 101 EPDM, 465, 640 EPDM [ethylene-propylenediene terpolymer, 587 EPDM rubber, 441 EPDM-based blends, 674 EPDM-ZnO-stearic acid systems, 443 epichlorohydrin, 263 epoxides, 378 epoxy, 261 epoxy aliphatic esters, 140 epoxy toluene oligomer, 264 EPR method, 88 EPR-signal, 88 equilibrium constants, 354 equilibrium contact angles, 526 error progression, 27 error tolerance, 1 ESR spectrum, 232 ester bond, 89 estimation of suspension viscosity, 707 ethanol, 207 ethanol content, 207 ethyl cellulose, 314 ethylene, 334 ethylene blocks, 334 ethylene diamine, 195, 536 ethylene-based polymers, 158 ethylenediamine, 235 ethylene-propylene copolymers, 81 ethylene-vinyl chloride copolymers, 335 EVA rubber blends, 473 excess flash, 453 excited-state complex, 401 excited-state quencher, 400 exnucleases, 102 extended chain segments, 128 extended polymer segments, 130 extended-chain crystals, 295 extensional deformation, 690 extent of blending, 146

866 exterior water-cooling, 66 extractables, 159 extraction, 220, 222, 420, 421 extraction of metal cations, 346 extraction process, 346 extraction solvent, 167 extruded strands, 630 extruder, 142 extrusion, 281, 456, 625, 626, 627 extrusion coating techniques, 579 extrusion compounding, 816 feed-back models, 4 feed-back networks, 4 feed-forward network architectures, 4 feed-forward networks, 4 feed-forward neural network, 2 fiber composites, 787 fiber glass, 84 fiber length, 819 fiber/nonwovens, 160 fiber orientation, 819 fiber volume fraction, 819 fiber-reinforced composite materials, 814 fiber-reinforced composites, 811, 833 fibers-matrices interface, 812 fibrillation behavior, 690 fibrillation (crazing), 292 fibrils, 310 fillers, 318 film, 318 film applications, 159 film extrusion, 278 film manufacturers, 159 film properties, 285 film size, 172 film-forming temperature, 172 filming, 172 films, 142, 405 fine structure of PET fibers, 839 flame retardant, 45, 426 flammability, 651 Fletcher nonlinear optimization, 20 Fletcher routine, 23 flexibility, 362 flexible materials, 761 flexible PVC, 137 flexible resin transfer molding, 817 flexural modulus, 164 floe formation, 130 flocculation, 130 flocculation activity of polymers, 71

Index flocculation process, 128 Flory-Huggin's interaction parameter, 637 Flory-Huggins theory, 18 flow behavior, 685, 701 fluid property prediction, 10 fluorescent poly (aryl ethers), 36 fluoroalkyl groups, 39 fluoroelastomer compound, 457 foaming, 769 folded-chain nucleation, 309 food-packaging applications, 579 fractional crystallization, 298 fractured molecules, 360 fragment length, 829 free bond rotation, 358 free energy, 304, 355, 371 free energy diagrams, 388 free energy of deflection, 355, 369 free energy of solution formation, 352 free energy of surface formation, 370 free radical initiators, 190, 198 free radical mechanism, 81 free radical polymerization, 243, 254 free radicals, 200, 244, 484 free-radical copolymerization, 182 free-radical scavenger, 401 freeze-thaw stable adhesive, 170 frequency, 393 FTIR-ATR, 218 function of emulsifier, 1% functional comonomer, 217 functional groups, 215, 218, 541 functionalized polyolefins, 164 functional materials, 411 functional monomers, 217, 230 fungal producers, 94 Fuoss relation, 63 gamma radiation, 119 gamma radiation grafting, 508 gamma radiation-induced grafting, 507 gamma radiation-induced preparation of water soluble polymers, 119 gas phase LLDPE, 157 gas sorption, 57 gauche bonds, 299 gel formation, 770 gel permeation chromatography, 288, 656 gelatin, 418 gel-chromatogram, 265 gel-chromatography, 264 genetic engineering, 413

Index Gibbs distribution, 365 glass fibers, 789 glass transition, 35, 54, 146, 476 glass transition temperatures, 275 glass-transition temperatures, 140, 571 glassy ionomers, 147 glassy state, 150 gloss, 651 glutaraldehyde, 218 golf balls, 651 Goodyear's Chemigum P83, 143 GPC curve of LLDPE, 158 GPC data, 425 graft copolymer, 164, 167, 204, 207, 417, 481, 496, 497, 725, 726 graft copolymerization, 481, 487, 491, 497, 501, 504, 541, 551, 796 graft copolymerization of cellulose, 529 grafted copolymers, 257 grafted polymer, 473 grafting, 255, 482, 483, 501, 511, 535 grafting by trapped radicals, 494 grafting by xanthate method, 532 grafting efficiency, 482 grafting in a nonwetting solvent, 538 grafting medium, 537 grafting of acrylic, 504 grafting of cellulose, 502, 529, 530, 533, 534 grafting of hydrophilic chains, 255 grafting of hydrophobic groups, 183 grafting of partially acetylated cellulose, 537 grafting of reactive antioxidants, 116 grafting process, 538 grafting rate, 530 grafting reaction, 116, 179, 501, 546, 551 grafting reaction of macromolecules, 546, 549, 550 graph theory, 26 graphical theory, 24, 27 gravimetric method, 510 green-flax fiber rovings, 807 grinding, 536 growth mode, 413 halogen groups, 255 halogenation reactions, 259 halosilanes, 45 hand-lay or contact molding, 816 Hansen solubility parameters, 206 hard metal cations, 348 hard PVC, 137 hardening, 137 hardness-impact resistance ratio, 163

head deflection temperature, 651 head-to-head structures, 324 head-to-tail polymer, 324 heat aging, 652 heat-aging resistance, 75 heat capacity, 28, 852 heat capacity of annealed PET fibers, 853 heat resistance of NBR, 555 heat stabilizers, 318 heat treatment of TPU, 139 heptane, 83 heteroatom, 373 heteroatomic dipoles, 393 heterocycles, 37 heterocyclic analogue, 39 heterocyclic system, 39 heterogeneous catalytic hydrogenation, 557 heterogeneous copolymerization, 69 hexagonal-orthorhombic phase transition, 304 hexyl group grafted polymers, 184 hidden neurons, 9 high inhibitive efficiency, 83 high melt tension LLDPEs, 158 high molecular weight esters, 23 high molecular weight polymers, 70 high performance polymers, 35 high temperature polymers, 35 high-density polyethylene (HDPE), 150 high-heat distortion grade blends, 659 high-impact strength, 640, 651 high-performance polymers, 424 high-pressure crystallization, 313 high-pressure crystallization kinetics, 308 high-pressure crystallization of PE, 295 high-pressure hexagonal phase, 301 high-pressure hexagonal phase of PE, 299 high-pressure mercury lamp, 379 high-pressure polyethylene, 73 high-processing temperatures, 445 high-shear dispersive melt mixing, 466 high-temperature stabilization, 81 high-tensile strength, 651 higher energy absorption, 712 higher-ordered structures, 98 hindered amine light stabilizers, 400 hindered phenols, 112 Hofmann degradation of poly aery lamide, 64 homeostasis, % homogeneous catalytic hydrogenation, 559 homopolymer, 145, 147, 150, 152, 254, 637 hot feed extrusion, 458 hot press molding, 818

868 HP LDPE, 281 HP LDPE film, 285 HPO-TA redox initiator, 170 HPO-TA system, 170, 173 Huggins equation, 63 humus, 417 hybrid composites, 805, 811 hydration of cations, 130 hydrocarbon polymers, 413 hydrocarbon thermal destruction, 83 hydroformylation of nitrile rubber, 567 hydrogen, 26 hydrogen atoms, 529 hydrogen bonding, 128, 207, 412 hydrogen bonds, 131, 302, 671 hydrogen peroxide, 533 hydrogen peroxide—reductant system, 489 hydrogenated carbon-carbon double bond, 558 hydrogenated nitrile rubber, 572 hydrogenation of liquid rubber, 566 hydrogenation of nitrile rubber, 556 hydrogenation of nitrile rubber latex, 566 hydrogenation process, 561 hydrohobic monomer, 217 hydrolysis, 100 hydrolysis of polyacrylamide, 64 hydrolysis processes, 777 hydrolysis reactions, 775 hydroperoxide, 110, 506 hydroperoxide groups, 318 hydrophilic microspheres, 199 hydrophobic chain, 200 hydrophobic groups, 179, 196 hydrophobic region, 179 hydrophobization, 185 hydrophobization of polyanionic polymers, 179 hydrophobized polymers, 180, 186, 187 hydroquinone inhibitor, 219 hydrorubeanic acid, 77 hydroxides of alkali metals, 83, 84, 87 hydroxy-containing polymers, 256 hydroxy 1 functional groups, 139 hydroxyl groups, 81 hydroxy propyl cellulose (HPC), 206 hyperfine splitting constant, 235 imaging technology, 432 imidazol rings, 260 immiscible polymer blends, 667 immiscible PP-NBR blends, 646 impact modifiers, 318, 761 impact performance, 151

Index impact strength, 164, 285, 292 impregnation of fibers, 7% in situ composites, 590 in situ formed copolymers, 635 incompatibility, 633 industrial applications, 771 industrial processes, 119 industrial production of detergents, 93 influence of ion content, 146 influence of plasticizers, 150 influence of radiation dose, 125 infrared spectra, 87 infrared spectroscopy, 330 ingredient weighing, 454 inhibition of polyethylene, 88 inhibitive efficiency, 86 inhibitors, 86, 510 initiation by free radicals, 529 initiation mechanism, 544 initiation-formation of macroradicals, 81 initiator, 169, 195, 200, 219, 379, 487 initiator concentration, 200, 204, 531 initiator end groups, 323 injection molded specimens, 630 injection molding, 460, 461, 623, 627, 816 injection moulding, 278 injection/transfer molding, 462 injection/transfer molding machine, 462 injection-molded, 582 inorganic residues, 119 insoluble polymer, 378 integrated model, 30 interchain crosslinking, 611 interface, 814 interface engineering, 714 interface modification, 823 interfacial bond strength determination, 828 interfacial processing and properties, 714 interfacial shear stress, 831 interlamellar tie chains, 284 intermeshing rotor, 455 intermolecular bonds, 365, 366, 368 intermolecular interaction, 367 internal bonds, 370 internal forces, 365 internal mixer rotor designs, 455 interparticle bridging mechanisms, 128 interpenetrating network polymer, 430 intramolecular synergism, 77 intrinsic viscosity, 268 inverse emulsions, 68 ion content, 145, 146, 148, 149, 152

Index

ion exchange resins, 260 ion neutralization, 441 ion-acetoacetanilide systems, 544 ion-alky 1 phenylcarbamate systems, 542 ion-exchange resin catalysts, 775 ion-scattering spectroscopy, 518 ionic aggregates, 444 ionic bonds, 146 ionic crosslinking, 146, 151 ionic emulsifiers, 197 ionic groups, 145, 148, 441 ionic size, 130 ionic thermplastic elastomer, 441 ionizing radiation, 65 ionomer blends, 145, 150 ionomer/homopolymer blends, 150 ionomer particles, 151 ionomers, 146, 441, 444 ionophores, 342 IR and NMR spectroscopy, 656 IR spectrum, 138 IR studies, 525 irradiation, 403 irradiation dose, 524 irradiation time, 403 IR-spectroscopy, 88 isobaric conditions, 296 isocyanate, 139 isolation of block copolymer, 759 isomerization, 89, 562 isomers, 317 isopropanol, 203, 207 isotactic polypropylene, 160, 161, 162 isothermal experimental density data, 14 isotropic melt, 384 ketones, 42, 262 kinetic chain lengths, 200 kinetic mechanism, 192 kinetic models, 193, 200, 205 kinetic parameters, 351 kinetic parameters of PS, 265 kinetic rate constant, 713 kinetic theory of emulsion polymerization, 192 kinetics of dispersion polymerization, 210 laminated composites, 811 lamination, 68 lateral methyl groups, 387 lateral substructure, 840 latex, 172 latex film, 169, 170

869 latex mixing, 653 latex particles, 205, 212, 213, 215 latex product, 189 LC polymers, 432 leather, 418 Levenberg-Marquardt method, 5 Lewis acid, 268, 271,483 Lewis acid systems, 266 Lewis catalysis, 264 light permeability, 73, 75 light screener, 399 light-emission processes, 400 lignin, 417 linear alcohol series, 206 linear alcohols, 206 linear chain, 362 linear polymer, 221, 277 li-polystyrene, 262 liposomal lipid bilayer, 181 liposomal membrane, 183 liposomes, 181 liquid carboxylated nitrile rubber, 565 liquid crystalline, 395 liquid crystalline polymer composite, 585, 587 liquid crystalline polymers, 42, 432, 585 liquid plasticizer, 137 liquid-crystalline polymers, 383 liquid-solid phases, 61 living polystyrene, 729 LLDPE, 156, 158, 281, 285 LLDPE film, 285 long-chain hydrocarbon, 419 long-chain hydrocarbon phenols, 419, 420 long-term stress crack resistance, 292 long-term thermoxidative stability, 112 low density polyethylene, 110 low ion contents, 147 low temperature brittleness, 284 low temperature flexibility, 571 low-density polyethylene, 426 low-density polyethylene (LDPE), 317 lower light permeability, 73 low-molecular compound reactions, 359 low-molecular models, 83 low-molecular weight azo compounds, 741 low-molecular weight oligomers, 190 low-pressure polyethylene, 73 low-temperature impact test, 142 lubricants, 137, 140, 141, 143, 318 macroazo initiators, 756 macro-azo-initiators, 739

870

macroazoinitiators, 728 macrobisperoxides, 726 macroinitiators, 726 macrointermediates, 725, 731 macromonomeric initiators, 730 macromonomers, 729 macroperoxy initiators, 758 macroperoxyazoinitiators, 728 macroperoxy initiators, 726 macroplasmodia, 97 macroporous particles, 220 macroporous structure, 222 macroradicals, 86 macroreticular cationic resins, 780, 782 Madaline III algorithm, 8 malate dehydrogenase reaction, 99 maleated EPDM rubber, 441 maleic anhydride, 182, 522 man-made fibers, 812 manufacturing process, 452 manufacturing technology, 578 manufacturing throughput time, 451 Mark-Houwink equations, 62 Marquardt algorithm, 5 Marquardt method, 6 masterbatch compositions, 466 material applications, 35 material for blends, 159 matrices, 814 matrix phase, 145 mechanical behavior, 50 mechanical loss, 145 mechanical properties, 110, 143, 146, 149, 150, 151, 160,391, 599, 630, 631, 848 mechanical properties of composites, 809 mechanical properties of ionomers, 145, 147 mechanical properties of natural fibers, 789 mechanical theory, 822 mechanism of polymerization, 424 medical purposes, 139 melt mixing, 654 melt mixing without vulcanization, 466 melt processability, 54 melt processing antioxidants, 112 melt viscosity, 152, 277 melted polymer, 140 melting, 298 melting behavior, 310 melting behavior of PE, 279 melting points, 47, 161 melting temperatures, 35, 311

Index membrane affinity, 183 mesogen, 383 mesomorphous phase, 843 mesomorphous phase in PET fibers, 844 mesophase, 385 mesophase-isotropic transition, 385 metabolite, 103 metal alkyl, 87 metal carbonyl complexes, 245 metal carboxylates, 331 metal cations, 342 metal chelates, 487 metal complex calxarenes, 339 metal complex system, 505 metal complexes, 250 metal deactivators, 109 metal dithiolates, 109 metal hydroxides, 145 metal ion, 128 metal oxide crosslinked carboxylic rubbers, 442 metallic reactions, 260 metallizing, 87 metallocene catalyst systems, 155, 156 metallocene catalysts, 162 metallocene isotactic poly propylenes, 161 metallocene structures, 161 metallocene-based LLDPE, 156 metallocene-based polyethylene processes, 157 metallocene-based polyethylene products, 158 metallocene-based polyolefins, 155 metallocenes, 155, 156 metallocene-trimethylaluminum system, 155 metallurgy, 137 metastabUity, 304 metastable folded chain, 2% methacrylate, 332 methacrylic acid, 504 methanol, 207 methlenes, 397 methods of grafting, 482 methyl groups, 387 methyl methacrylate, 738 methyl methacrylate monomers, 250 methylene chloride, 149 methylol derivatives of polyacrylamide, 71 methylol groups, 797 micelle concentration, 192, 1% micelle structure, 196 micelles, 197 microdebond test, 831 microfibrillar substructure, 849 micron-size range, 211

Index microphase separations, 601 microscopy, 655 microspheres, 603 microstructure, 841 migration of antioxidants, 115 migration of plasticizers, 137 mineral fibers, 813 miscibility, 14 mixed solvents, 175 mixed solvent-water systems, 538 mixing characteristics, 475 mixing duty, 717 mixing equipment, 624 mixing step, 441 MM A polymerization, 246 mobile random conformation, 296 model calculations, 100 modification by blending, 526 modification by chemical methods, 519 modification by corona treatment, 526 modification by flame treatment, 527 modification by gas plasma treatment, 527 modification of natural polymers, 415 modifications of Levenberg-Marquardt method, 6 modified nitrile rubber, 568 modified polyolefin, 471 modifiers, 137 modulus, 145, 147, 162, 285, 391, 819 modulus of dispersed TLCP fibers, 698 modulus/rigidity, 651 molar volume, 28 mold, 460 mold cavity, 163 mold filling, 816 molding, 318, 458 molecular architecture of homogeneous polymer molecules, 155 molecular calculations, 98 molecular mass, 85, 97 molecular mechanism, 85 molecular parameters, 657 molecular structure, 277 molecular weight, 68, 84, 111, 120, 140, 167, 185, 187, 190,511 molecular weight distribution, 156,277,287,290,511 molecular weights, 216 molecular-chain scission, 88 monoamic acid of TAPO, 46 monocyclopentadienyl-amide derivatives of titanium complexes,, 156 monofunctional unit, 10

871

monomer, 68, 197, 198, 203, 216, 377, 670, 741 monomer concentration, 510, 530 monomer diffusion, 199 monomer droplets, 192 monomer phase, 221 monomer swelling step, 212 monomer system, 509 monomer type, 208 monomer-benzene solutions, 248 monomer-soluble initiators, 201 monomer-to-water ratio, 198 monomeric analog, 347 monomeric ethylene, 82 Monte Carlo method, 301 morphological parameters, 658 morphological studies, 644 morphologies of the PP-LCP melt blends, 626 morphology, 155, 219, 284, 310, 586, 623, 625, 669 Morton equation, 212, 214 mosiac grafting, 494 mould cavity, 289 mouse macrophages, 179 multicavity injection molds, 462 multimodal assemblies, 710 multiphase rheology, 707 multistage emulsion polymerization, 212 narrow composition distribution, 160 narrow tacticity distribution, 160 Na-SPS ionomer, 148 National Science Foundation, 29 natural fiber mat-reinforced thermoplastics (NMT), 804 natural fiber modification, 795 natural fiber reinforced plastics, 787, 799 natural fiber reinforced thermosets, 802 natural fiber-reinforced thermoplastics, 803 natural fibers, 787, 791, 795, 813 natural monomers, 418, 435 natural polymers, 412, 413 natural rubber, 411, 417, 465 natural rubber-based blends, 675 naturally occurring monomers, 434 NBR-based blends, 676 NBR-PVC blends, 472 network points, 364 network training data, 5 neural network architecture, 3 neural network performance, 15 neural network training processes, 11 neural networks, 1, 4, 5 neuron, 3,7

872

neurotoxin, 61 neutralization, 146, 148 neutralizing agents, 145 neutron activation analysis, 249 n-hexane, 172 nitrile rubber, 555 nitrile rubber technology, 574 nitriles, 42 nitrocellulose based printing ink, 527 nitrogen, 83 nitrogen ylide, 374, 376 NMR-spectra, 98 NMR spectroscopies, 569 NMR spectroscopy, 310 NMR techniques, 390 noncatalytic hydrogenation, 567 noncovalent binding of small cations, 99 non-Gaussian coils, 63 non-homopolymerizable antioxidants, 116 nonlinear optimization training, 5 nonpolar groups, 328, 333 nonpolar plasticizer, 150 nonrandom chain scission, 100 normal polyethylene, 76 nucleation, 202, 217, 305, 308 nucleation of ECC, 305 nucleation period, 208 nucleation processes, 308 nucleic acids, 413 nucleophilic displacement, 227 nucleophilic property, 37 number of passes, 456 OH radical, 493 oil production, 71 oil recovery, 119 oil-producing industry, 71 oligomer, 217 oligomer chains, 208 oligomer compounds, 201 oligomer units, 190 oligomeric material, 212 oligomerization, 427 oligomers, 208 oligoperoxides, 727 one-dimensional growth, 309 one-step processing, 142 optical birefringence, 2% optical material, 73 optical properties, 655, 855 optimization procedure, 706 optimized mixing line, 456

Index

organic hydroperoxide-amine systems, 231 organic metal salts, 326 organic peroxide-aromatic tertiary amine system, 227 organic phase, 216 organic phase-seed latex ratios, 220 organic phosphites, 83 organic solvents, 50, 68, 98, 175 organometallic compounds, 243 organometallic photoinitiators, 244 orientation, 291, 793, 844, 847 orientation function, 848 orthorhombic crystals, 304 orthorhombic lattices, 311 orthorhombic phase, 301, 311 outdoor weathering, 110 oxidation, 261 oxidation of organic materials, 105 oxidation process, 105 oxidation-reduction reaction, 195 oxidative degradation, 110 oxidative deterioration, 109 oxidising, 109 oxidizer, 706 oxidizing agents, 106 oxime groups, 344 oxygen, 115 oxygen pressure, 106 oxygen-free aqueous solution of acrylamide, 66 ozone degradation of polyacrylamide, 63 packing arrangement, 386 palladium catalyst, 558, 564 Pandex series of TPUs, 143 panel products, 580 paper, 119 paraffin, 173 particle fluid assemblies, 708 particle formation, 204 particle size, 200 particle size effect, 778 particulate composites, 811 patterns of crystallization, 295, 303 pemcillium cyclopium, 93 performance enhancement, 707 performance parameters, 706 permeability, 57 permeability to organic solvents, 172 peroxidation method, 509 peroxidation reactions, 491 peroxide, 227 peroxide cured system, 469

Index

peroxide decomposers, 112, 401 peroxide initiator, 229 peroxide vulcanization of polyethylene, 90 peroxide-amine initiation systems, 230 peroxide-functional monomer, 229 peroxy, 81 peroxy acid redox systems, 506 peroxy group, 749 peroxydiphosphate system, 486 peroxydisulfate initiators, 195 peroxydisulfate system, 486 Persian Gulf countries, 411 persulfate-amine systems, 232 persulphate systems, 505 PET, 318 PET- and PBT-based blends, 673 PET fibers, 839, 848 pH moisture, 30 pH values, 130 phase behavior, 383, 389 phase boundaries, 302 phase diagram of PE, 299 phase diagrams, 310 phase transfer catalysts, 42 phase transition, 312 phase-separation, 334 phase-value ratio, 658 PHB, 103 phenol/thiophenol based polymer precursors, 37 phenols, 87, 117 phenoxide, 42 phenyl group grafted polymer, 184 phenyl phosphine oxide, 39 phosphonium ylide, 377 phosphorus, 427 phosphorus containing polyimides, 46 phosphorus containing polymers, 45 phosphorus-containing stabilizers, 89 phosphorylation, 426 photoantioxidant activity, 113 photoantioxidants, 113 photochemical chain lengthening process, 243 photochemical postpolymerization, 243 photocuring, 257 photodegradation, 493 photo-induced charge transfer, 236 photo-induced charge-transfer initiation, 541 photo-induced charge-transfer initiation systems, 227 photo-induced method, 506 photoinitiated free radical polymerization, 256 photoinitiated grafting, 253

873

photoinitiating systems, 250 photoinitiation, 243, 247, 248 photoinitiation of polymerization, 244 photoinitiation processes, 244 photoinitiators, 245, 248 photoirradiation, 236 photolysis, 247 photolytic process, 251 photo-oxidative degradation, 404 photo-oxidative thermal degradation, 399 photopolymerization, 236, 243, 379 photoreaction, 249 photosensitivity, 266 photosensitizer, 114 photosensitizing, 113 photostability of polymers, 112 photostabilization of polyolefins, 112 phthalic anhydride, 83 physical activation, 483, 536 physical modification method, 259 physical properties, 110, 446, 848 physical properties of hydrogenated nitrile rubbers, 570 physical property predictions, 1 physical requirements, 112 physical structure of natural fibers, 793 physico-chemical properties, 97 physico-mechanical properties of polystyrene, 269 piezochromism, 176, 177, 178 pigments, 318 piperazine, 43 Planck constant, 308 planetary mixer, 717 plasma grafting, 495 plasma treatment, 825 Plasmodium, 97 plastic, 465 plastic blend, 466 plastic industry, 582 plastic materials, 93 plasticating, 141 plasticating extruder vs Kokneader, 719 plastic-cellulose networks, 577 plasticization of vinyl chloride, 335 plasticization stages, 816 plasticizer, 12, 14, 137, 139, 140, 141, 142, 143, 150 plasticizer leeching, 137 plate-assay, 99 PMMA grafted, 240 PMMA ionomers, 150 PMMA-based ionomer, 145, 147, 149, 150 polar comonomer, 216

874

polar component, 522 polar dimethylformamide (DMF), 148 polar groups, 328, 331 polar monomers, 193, 1% polar solvent, 149 polarization microscopy, 73 polidispersity, 203 poloxamers, 765, 771 poly(cyclohexane oxide), 748 poly(heteroarylene) sulfide, 39 poly(malic acid), 93 poly(methyl methacrylate) (PMMA) ionomer, 145 poly(methylmethacrylate), 332 poly (sty rene), 743 poly(styrene-acrylamide), 216 polyvinyl acetate) (PVAc) latexes, 167 polyvinyl alcohol) (PVA), 167 poly(vinyltrichloacetate), 254 polyacrylamide, 62, 64, 66 polyacrylamide forms, 130 polyacrylate, 458 polyacrylic acid, 203, 218 polyaddition, 737 polyaddition reactions, 741 polyamides, 69 polyamine disulfides, 79 polyamine disulphides, 90 poly amines, 218 polyanionic polymers, 179, 184, 187 polybibenzoates, 384, 391, 393, 394 polychlorotrifluoroethylene, 297 polycondensation, 47, 738, 748 poly disperse particle processing, 708 polydispersity indices, 221 polyelectrolytes, 119, 128 polyester-based TPU, 143 polyesters, 93, 103, 143, 757 polyether-elastomer block copolymers, 634 polyetherimides, 54 poly ethers, 43, 50 polyethylene, 76, 86, 89, 137, 155, 277, 485 polyethylene macroradical, 87, 88 polyethylene molecules, 87 polyethylene oxide, 765 polyethylene terephthalate, 297 polyethylene terephthalate (PET) fibers, 839 polyethylene thermal destruction, 87, 88 polyethylene-based ionomers, 148 polyethylene-henequene-sand laminates, 834 polyimides, 47 polymalatase, 94, 101, 102 polymer activities, 20

Index polymer alloys, 137, 601 polymer backbone, 35, 334 polymer blends, 138, 143, 290, 611, 633, 649, 685 polymer branching, 126 polymer chains, 129, 155, 191, 200, 485 polymer chemistry, 318 polymer crystalline morphology, 280 polymer drugs, 187 polymer educators, 649 polymer efficiency, 130 polymer equilibrium, 365 polymer market, 61, 164 polymer microspheres, 218 polymer morphology, 710 polymer networks, 252 polymer particles, 201 polymer precursors, 37 polymer property predictions, 24 polymer salt, 94 polymer solutions, 99 polymer stabilization, 117 polymer synthesis, 54, 156 polymer synthetic activity, 97 polymer systems, 19 polymer volume fraction, 20 polymer-bound stabilizers, 402 polymeric blends, 137, 143 polymeric calxarenes, 339, 346 polymeric chains, 126 polymeric flocculants, 130 polymeric materials, 160, 399, 599 polymeric matrix materials, 815 polymeric modifiers of PVC, 143 polymeric plasticizer for PVC, 138 polymeric salts, 374 polymeric UV absorbers, 402 polymeric ylides, 374 polymerizable monomer, 481 polymerization, 65, 168, 486, 817 polymerization by ylides, 373 polymerization conditions, 200 polymerization equipments, 193 polymerization medium, 201 polymerization methods, 168, 189 polymerization of acrylamide, 66 polymerization of liquid propylene, 161 polymerization of styrene, 376 polymerization of VAc, 170 polymerization of vinyl chloride, 319 polymerization procedure, 193 polymerization rate, 66, 193, 199, 208, 209 polymerization reaction, 35, 191, 193

Index

polymerization temperature, 193, 196, 199,201, polymerization times, 212 polymerization variables, 200 polymerization volume, 192 polymethylenic sequences, 394 polyolefin materials, 155, 164 polyolefin stabilizers, 77 polyolefin thermal destruction, 85 polyolefin-polyamide blends, 668 polyolefins, 69, 73, 81, 112, 164, 471 polyphosphates, 46 polypropylene (PP), 317 polypropylene, 81, 82, 160 polypropylene block, 162, 164 polypropylene oxide, 765 polypropylene structures, 159 polypyrocatechin borate, 88, 89 polystyrene, 85, 218, 221, 259, 263, 345 polystyrene microspheres, 218 polystyrene (PS), 317 polystyrene seed emulsion, 218 polystyrene seed particles, 218 polystyrene-polyhydroxyethyl-methacrylate (PS/ PHEMA), 217 polyurethane, 144, 255 polyurethane elastomers, 144 polyurethane-grafted copolymer, 256 polyvinyl alcohol, 69, 222, 295, 503 polyvinyl chloride, 69 polyvinylchloride, 137 pore blocking, 71 pore size distribution, 221 pore structure, 221 pore volumes, 221 porous film, 164, 167, 171, 172, 174, 178 porous materials, 170 porous polyvinyl alcohol)-poly (vinyl acetate) composites, 167 porous PVA film, 172 porous PVA-PVAc composite films, 175 porous PVA-PVAc composites, 172 porous uniform latex particles, 189 potassium acetate, 329 potassium hydroxide, 83 potassium peroxydisulfate, 200, 219 potassium phosphate buffer, 101 precursor block copolymers, 602 precursor polymer, 204 preirradiation method, 509 pre-irradiation method, 491 preparation of PBAs, 653 press controls, 459

875

pressure-crystallized powder, 310 preventive antioxidants, 109 primary radicals, 190 processability, 157, 280, 288, 290, 651 processability behaviors, 442 processing, 625 processing advantages of metallocenebased polyolefins, 164 processing aids, 318 processing conditions, 147, 287, 589 processing difficulties, 283 processing equipment, 716 processing industry, 291 processing limitations, 290 processing machine, 140 processing methods, 143 processing of natural fiber reinforced plastics, 802 processing properties, 141, 470 processing techniques, 715, 816 processing temperature, 140, 141, 477 processing time, 142 progressive insolubilization, 64 properties of dry polymer, 62 properties of ionomers, 148, 152 properties of polypropylene, 161 properties of the PVC/TPU blends, 142 properties of unvulcanized elastomer-plastic blends, 473 property enhancement, 152 property relationships, 156 propylene oxide, 766 proteins, 413 proton transfer, 240 PS-based blends, 671 pseudoplastic fluids, 445 PTE, 17 pulpal irritation, 229 pultrusion, 818 puncture resistance, 284, 285 purification, 45 purification methods, 768 PVA, 167, 170, 174 PVA cells, 173 PVA concentration, 171 PVA film, 174 PVA-PVAc composite film, 176 PVAc grafted on PVA, 175 PVAc latex, 169, 172 PVAc latexes, 167, 168, 169, 170, 173 PVAc latex films, 167 PVAc particles, 169, 171, 177 PVC, 138, 142

876

PVC compositions, 317 PVC granules, 142 PVC growth, 317 PVC materials, 137 PVC matrix, 328 PVC modified by TPU, 143 PVC plastics, 137, 144 PVC products, 137 PVC resins, 140, 141, 142 PVC stability, 327 PVC stabilization, 322, 326 PVC/TPU/ABS blend, 144 PVC/TPU blends, 141 PVGTTPU polymeric blends, 144 PVC/TPU system, 141 pyrazole group, 37 pyridazine, 39 quaternary salts, 374 quenched polypropylene, 73, 74 quenchers, 402 quenching, 400 quinoxaline ring, 37 radial lip shaft seal, 451 radiation, 120, 126, 399 radiation chemistry, 56 radiation crosslinking of PS, 147 radiation grafting, 509 radiation resistance, 56 radiation sensitivity, 509 radiation-induced graft copolymerization, 489 radical catalysis, 259 radical cation transfer, 748 radical chain reactions, 749 radical copolymerization, 69 radical initiator, 483, 484 radical polymerization, 81 radical trapping, 232 radically polymerizable monomers, 430 radical-radical transfer, 744, 749 radicals, 81, 110,545 Raman spectroscopic studies, 300 random copolymer, 481 random initiation, 322 rate constants, 210 rate of flocculation, 128 rate of formation, 244 rate of polymerization, 120, 238 rat intestinal cells, 180 raw materials, 141, 452 reaction kinetics, 561

Index reaction rate, 64 reaction temperature, 531 reaction time, 530 reactive blends preparations, 466 reactive compatibilization, 667, 668 reactive compatibilization technique, 646 reactive compatibilizing agents, 671 reactive injection molding (RIM), 722 reactive processing, 710 reactive processing methods, 116 reactively processed blends, 468 reactively processed polymer blends, 465 reactivity of maleic anhydride grafted PP, 474 reciprocating screw machines, 460, 461 recrystallization rates, 730 recycles plastic, 578 redox,227 redox initiation systems, 227 redox initiators, 195 redox method, 503 redox system, 485, 487 redox-initiated polymerization of vinyl monomers, 227 redox-initiated polymerization rate equation, 230 reducing agent, 106, 195 refractive index, 174, 175, 177 refractive index of materials, 174 regenerated fibers, 812 regular copolymer, 481 reinforced reaction injection molding (RRIM), 818 reinforcement, 812 reinforcing, 152 remelting, 816 repolymerization of monomer, 220 repolymerization step, 220 reprocessability studies, 445 residual lignin in pulp, 531 resin matrix, 819 resin processability, 285 resin transfer molding, 817 reversed-phase chromatography, 101 rheological determination, 690 rheological parameters, 613, 615, 616, 619, 621 rheological properties, 163 rheology, 280, 585, 630, 669 rhodium catalyst, 559 Rhus vernicefera, 420 rigidity, 360, 362 RIM energy composites, 722 ring structures, 155 roll speeds, 456 roll temperature, 456

Index

ROOH-amine systems, 231 rotary isometry, 353 rotating speed of the screw, 142 rotomoulding, 278 rubber matrix, 476 rubber-plastic blends, 465, 634 rubbery plateau, 146 rubbery polymers, 420 rubbery region, 146 rubbery type ionomer, 152 ruthenium complexes, 562 saccharide molecule, 252 sample thickness, 73, 75 saponification of acrylonitrile, 62 scanning electron microscope, 511 scanning electron microscopy, 497 scanning electron microscopy (SEM), 611 scattering process, 73 Scott-Gilead, 114 screeners, 402 screw/ram machine, 461 SEBS rubber-based blends, 675 secondary ion mass spectroscopy, 518 secondary nucleation, 203, 208 seed latex dispersion, 218 seed particles, 213, 215, 220 seed polymer, 213 seeded emulsion polymerization of styrene, 219 selection of emulsifiers, 197 selection of TPU, 139 selective extraction experiments, 345 self- or aggregative nucleation mechanism, 201 self-extinguishing, 46 semi-crystalline character, 839 SEM micrographs, 208, 592 shaft seal, 453 shaft seal molding process, 459 shear flow of Newtonian fluids, 586 shear history, 281 shear modification, 281 shear modulus, 618 shear stress, 445 shear thinning behavior, 281 sheet, 318 shellac, 418 short chain branching level, 157 short fiber-reinforced plastic composites, 811 short-wavelength ultraviolet (UV) rays, 399 shrinking, 137 side chains, 81 sigmoid, 3

877

silica filler, 441, 445 silicone elastomers, 27, 31 silicone structure, 31 silicone-EPDM blend, 526 silk, 411,418 siloxanes, 10 silyl groups, 42 silyl method, 45 silylated monomers, 42 silylation, 45 simple tension, 147 single cavity mold, 461 single drop mill, 456 single fiber pull out test, 830 single-stage soapless emulsion polymerization, 215 size distribution, 204, 216 size exclusion chromatography, 94 skin layer, 444 slurry phase processes, 155 small-angle neutron scattering (SANS), 145 small-angle x-ray scattering (SAXS), 145 SMC molded plastics, 802 smectic mesophase, 388 Smith-Ewart theory, 190, 192, 197 soapless emulsion polymerization, 199 soapless emulsion process, 200 sodium bicarbonate, 196, 221 sodium lauryl sulfate, 221, 222 sodium metabisulphite, 195 sodium nitrile, 63 sodium sulfide, 37 sodium thiocyanate, 341 softening points, 47 soil structure formation, 71 solidification, 284 solubilities of acrylamide, 61 solubility, 28 solubility method, 655 solubilization, 770 solubilization of aromatic polyimines, 50 solubilized emulsifier molecules, 190 solubilized monomer molecules, 190 solvatochromism, 175 solvent activity, 17, 24 solvent extraction mechanism, 344 solvent resistance, 651 solvent solution polymerization, 67 specific gravity, 652 specific molar volumes, 12 specific surface area, 221 specific volumes, 11 specific volumes of siloxanes, 10

878

Index

spectral properties, 98 spectroscopic measurements, 4% spectroscopic methods, 767 spectroscopic techniques, 517 spectroscopy analysis, 10 spherulite, 75 spherulite transparent structure, 76 spontaneous dissociation, 100 spontaneous ester hydrolysis, 100 spray molding, 817 stability, 140, 320 stability of resins, 777 stabilization, 39, 83, 399 stabilization of polyolefins, 105 stabilization of polymers, 399 stabilization of PVC, 318, 327 stabilization technology, 105 stabilizer, 117, 137, 140, 143, 204, 205, 206, 222, 300, 402, 404 stabilizer concentration, 79, 205 stabilizing, 112 stabilizing action, 77 stabilizing activity, 79 stabilizing effect, 91 standard error-back-propagation, 7 star block copolymers, 601 starch, 413 static vulcanization, 467 steady-state conditions, 307 stereoregular polypropylene, 159 stereospecific catalyst, 277 steric barrier of dissolved polymer, 189 steric hindrance, 777 steric stabilizer, 201 steric stabilizer systems, 203 stiffness, 145, 150, 292, 642 stirred polymerization reactor, 204 stirring rate, 780 stock blender, 456 storage modulus, 145, 393 strain-to-fracture, 151 strength, 642, 819 stress, 393, 592 stress distribution, 832 stress-strain plots, 391 structural backbone, 415 structural control, 726 structural defects in PVC, 319 structural effects of amine, 231 structural information, 97 structure-property relationship of graft copolymers, 641

structures of functional monomers, 230 styrene, 164, 198, 199, 217, 218, 222, 250, 375, 750 styrene monomer, 216 styrene-divinylbenzene, 213, 215 styrene-methylmethacrylate mixture, 222 substitution of chlorines, 331 successive grafting, 494 sulfate radicals, 195 sulfide, 57 sulfide linkages, 35 sulfonated EPDM rubber, 441 sulfonated ethylene-propylenediene terpolymer, 152 sulfone, 41 sulfone groups, 57 sulfoxide, 57 sulfur, 37 sulphonium ylide, 378 sulphur, 83 sulphur compounds, 109 supercooling, 303, 304 superoxide production, 179 superstructure, 839 surface activity, 715 surface coatings, 419 surface free energy, 815 surface modification, 795 surface modification of fillers, 825 surface modification of polymers, 823 surface modifiers, 762 surface properties, 517 surface properties of modified polymers, 517, 519 suspension (co)polymerization, 67 suspension polymerization, 67 suspension polymerization techniques, 189 swelling, 215, 220 swelling agent, 212, 213 swelling degree of latex films, 170 swelling measurements, 511 swelling process, 220 swelling properties, 742 swelling steps, 189 swollen emulsion polymerization, 189 swollen particles, 213 synchrotron experiments, 385 syndiotactic arrangement, 325 syndiotactic content, 325 syndiotactic polypropylene, 163 synergism, 117 synthesis of macroinitiators, 756 synthesis of the microspheres, 602 synthesis of TPU, 139 synthesizing TPU, 139

Index

synthesizing block copolymers, 755 synthetic fibers, 416, 813 synthetic polymers, 102, 417, 481 synthetic structures, 412 synthetic wood, 577, 582 tacticity, 325 tailor polymer performance, 157 take-off system, 454, 455 tangential rotor, 455 tear strength, 285, 291 technical applications, 787 technical applications of natural fibers, 787 technology of compounding, 649 TEM microspheres, 603 TEM scans, 149 temperature dependencies, 363 temperature variations, 68 tensile, 692 tensile modulus, 52 tensile properties of metallocene, 158 tensile strength, 110, 145, 147, 148, 150, 151, 285, 594, 850 tensile strength versus blend ratio, 151 tensile stresses, 587 tension, 146 terminating agents, 741 termination mechanism, 750 termination reaction, 192 tertiary amine, 234 tertiary butyl hydrogen peroxide (TBHP), 484 tertiary chlorine, 323 tetrafluoroethylene (TFE), 246 textiles, 119 texture diffractogram, 846 TGA curves, 271 theory and mechanics of reinforcement, 831 theory of compatibility, 138 thermal analysis, 4%, 511 thermal and thermomechanical analysis, 655 thermal decomposition, 195 thermal decomposition of polyethylene, 81 thermal degradation of PVC, 320 thermal dehydrochlorination, 327 thermal destruction, 84, 87, 88, 89, 90 thermal destruction of polyethylene, 81 thermal destruction of polypropylene, 84 thermal initiators, 195 thermal motion, 85 thermal oxidative breakdown, 77 thermal oxidative breakdown of polyethylene, 77 thermal properties, 851

879 thermal stabilities, 330 thermal stability, 52, 332, 334, 651 thermal stability of polyethylene, 81 thermal stability of PVC, 318, 331 thermal stabilization, 111 thermal stabilization of polyolefins, 111 thermal treatment, 148, 149 thermal treatment of PVC, 326 thermal volume expansivity, 851 thermochromism, 176 thermodynamic modelings, 300 thermodynamic parameters, 390 thermodynamic property estimations, 4 thermodynamic terms, 300 thermodynamics of binary TLCP/ thermoplastic blends, 585 thermodynamics of nitrile rubber hydrogenation, 556 thermogravimetric analysis, 318 thermooxidation, 351 thermoplastic behavior, 449 thermoplastic composites, 577 thermoplastic elastomer blends, 633, 634 thermoplastic elastomers, 647, 653 thermoplastic fibers, 577 thermoplastic market, 155 thermoplastic polyurethane, 137 thermoplastics, 416, 583 thermostability, 271 thermostabilizing efficiency of poly amine disulphides, 91 thermostatic bath, 194 thermotropic main-chain liquid crystalline polymers, 623 thermotropic polybibenzoates, 383 thiazole groups, 41 thin films, 149 thiocarbomoyl ligand, 114 thioether linkages, 52 thiol-terminated oligopropylene, 164 thiophenoxide, 42 thiopropionate esters, 109 thiourea, 63 three-dimensional crystallinity, 395 three-dimensional structures, 604 three-dimensionally ordered crystals, 383 throughput time, 464 titration methods, 191 TLCP blends, 591 TLCP phase, 595, 597 tough composites, 623 toughness, 147, 148

880

toughness of PVC, 139 toxic monomer, 42 toxicity, 102

TPU, 137, 138, 142 TPU modified by PVC, 143 training patterns, 9 transesterification, 480 transfer agent, 735 transfer of hydrogen, 88 transformation, 416 transformation to porous film, 170 transition bond energy, 128 transition metal atom, 155 transition metal chelates, 247 transition rates, 309 transition temperature, 314 transmission electron microscope (TEM), 146 transmittance spectra, 174 transmitted light flux, 73 transparency, 110, 140, 163 transparent films, 52 transparent heterogeneous blends, 659 transparent poly olefins, 73 transparent random copolymers, 162 transport properties, 57 transversal modulus gradient, 716 treatment with isocyanates, 797 triclinic crystallographic system, 841 triglyceride esters, 15 triisopropyl phosphine, 562 triple point, 310 tubular blown film, 278 tubular film blowing operation, 283 twin-screw extruder, 142, 455 twin-screw extrusion, 721 two-mill take-off system, 456 two-step polymerization technique, 218 two-step processing exruders, 142 two-wing tangential rotor mixers, 455 type of counterion, 149 types of hexagonal phases, 301 ultrasonic waves, 65 ultraviolet, 112 ultraviolet irradiation, 75, 76, 824 ultraviolet radiation, 65 ultraviolet spectra, 379 ultraviolet (UV) light, 244 ultraviolet (UV) spectra, 320 unfills, 453 uniaxial stretching, 283 uniform latex particles, 189, 191, 193, 211

Index unsaturated aromatic acids, 262 unsaturated end groups, 323 unsaturated esters, 42 unsaturated polyester-styrene mixture, 257 unsaturation, 319, 560 unstabilized PVC, 318 uranium, 342 urushiol, 421 use of graft copolymers, 497 UV absorber, 109, 114, 400, 402, 404 UV analysis, 232 UV stabilizers, 113 UV-stabilizer, 115 UVA activity, 117 valancy of cations, 130 vanadium V(V) systems, 504 van der Waals force, 137 van der Waals radii, 358 van der Waals volume, 28 VCZ cation radical, 236 VCZ fluorescence, 236 vinyl chloride, 317, 321 vinyl monomers, 236, 237, 253, 485, 501, 541 vinyl polymerization, 227, 244, 246, 731, 732 vinylchloride monomers, 141 viscoetastic composite, 705 viscoelastic phenomena, 392 viscoelastic properties, 649 viscosities of the siloxanes, 12 viscosity, 12, 141, 169, 445, 475, 685 viscosity predictions, 13 viscous flow temperature, 141 viscous properties, 63 voids, 821 vulcanization, 74 water content, 207 water solubilities, 216 water treatment, 127 water-bearing stratum, 71 water-in-oil emulsions, 71 water-soluble inhibitor, 213, 220, 222, water-soluble initiator, 216 water-soluble polymers, 119, 130 wavelength, 73, 178 weak boundary layers, 795 weathering data, 31 weathering events, 31 weight-average molecular weight, 130 wet adhesive bonding process, 579 wettability, 815

Index wetting, 537, 770 wetting agents, 771 wood,411 wood-polymer composites, 581 wood-polypropylene fiber composite, 578 wool, 418 xanthate method, 506 x-ray diffraction, 389 x-ray diffractograms, 389 x-ray fluorescence spectroscopy, 517 x-ray measurements, 300

x-ray photoelectron spectroscopy, 218, 519 x-rays diffraction method, 656 yellowing, 110 ylides, 373, 374, 379 Ziegler-Natta catalyst, 161, 164, 277 Ziegler-Natta polymeric products, 157 zinc stearate, 442, 444 zinc stearate additive, 152 Zn-salts of a butyl acrylate/acrylic acid polymer, 147

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