Short Fibre-Polymer Composites

Short fibre-polymer composites

Short fibre-polymer composites Edited by

S K DE and J R WHITE

WOODHEAD PUBLISHING LIMITED CAMBRIDGE ENGLAND

Published by Woodhead Publishing Limited, Abington Hall, Abington, Cambridge CBl 6AH, England First published 1996, Woodhead Publishing Limited

0 1996, Woodhead Publishing Limited Conditions of sale All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopy, recording, or any information storage and retrieval system, without permission in writing from the publisher. While a great deal of care has been taken to provide accurate and current information, neither the authors, nor the publisher, nor anyone else associated with this publication, shall be liable for any loss, damage or liability directly or indirectly caused, or alleged to be caused, by this book. British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library. ISBN 1 85573 220 3

Cover design by the ColourStudio Typeset by Vision Typesetting, Manchester, England Printed by Galliard (Printers) Ltd, Great Yarmouth, England

Contents

Preface Contributors 1 Survey of short fibre-polymer composites

ix xi 1

J R WHITE AND S K DE

1.1 1.2 1.3 1.4 1.5 1.6 1.7

2

Introduction Composition Morphology of short fibre reinforced polymers Mechanics of short fibre reinforced polymers Measurement of fibre orientation distribution Properties of fibre reinforced polymers Conclusions References

Short fibre filled thermoplastics

1 2 5 7 14 17 18 19

21

J R WHITE

Introduction Materials Fabrication - injection moulding Properties Effect of fabrication on properties Conclusions References

21

3 Thermosetting short fibre reinforced composites

54

2.1

2.2 2.3 2.4 2.5

2.6

21 24 35 44 50 50

S B WlLKlNSON AND J R WHITE

3.1 Introduction 3.2 Curing characteristics 3.3 Thermosetting resin types

54 54 56

Contents

vi

3.4 3.5 3.6 3.7 3.8 3.9 3.10

Fabrication methods Reinforcing fibres Fibre orientation Fillers and other additives to thermosets Residual stresses Properties of short fibre reinforced thermosets Conclusions References

4 Short fibre-thermoplastic elastomer composites

61 64 66 70 73 77 81 81 84

G B N A N D 0 A N D B R GUPTA

4.1 4.2 4.3 4.4

Introduction Classification of TPEs Short fibre-elastomer composites Parameters influencing the characteristics of short fibre-polymer composites 4.5 Short fibre-thermoplastic elastomer composites (SF-TPE) References

5 Composites of polychloroprene rubber with short fibres of poly(ethy1ene terephthalate) and nylon

84 85 87 89 96 113

116

M ASHIDA

5.1 5.2 5.3 5.4 5.5

Introduction Preparation of composites Mechanical and viscoelastic properties Effect of absorbed water Dynamic fatigue of composites Acknowledgement References

6 Properties and processing of short metal fibre filled polymer composites

116 116 117 131 135 142 142

144

D M BlGG

6.1 6.2 6.3 6.4 6.5 6.6 6.7 6.8

Introduction Metal fillers Polymers Electrical properties Thermal properties Mechanical properties Processing Effect of environment on properties References

144 146 149 149 156 156 162 165 166

Contents

7 Electrically conductive rubber and plastic composites with carbon particles or conductive fibres

vii

168

P B J A N A , A K MALLICK A N D S K DE

7.1 7.2 7.3 7.4 7.5 7.6 7.7 7.8 7.9 7.10 7.11

Introduction Percolation phenomena in conductive polymer composites Mechanism of electrical conduction Effects of processing factors on electrical properties Effects of polymer matrices on conductive network formation Effects of the types of filler, their geometry and morphology Effect of structural deformation on electrical resistivity Hydrostatic pressure effect on resistivity Temperature effects on volume resistivity Galvanomagnetic properties of conductive rubber composites EM1 shielding effectiveness References

8 Design and applications of short fibre reinforced rubbers

168 168 170 172 175 175 176 177 177 179 182 188 192

M C H LEE

8.1 8.2 8.3 8.4 8.5 8.6 8.7 8.8 8.9

Introduction Materials Sample preparation Test procedure Stress-strain behaviour of unidirectional short fibre reinforced elastomer composites Fracture morphology, fibre configuration and deformation mechanism Stress-strain equation for unidirectional short fibre reinforced elastomer composites Modulus properties of unidirectional short fibre reinforced elastomer composites and applications Concluding remarks References

9 Design considerations and end-use applications of short fibre filled rubbers and thermoplastic elastomers

192 193 193 194 195 196 202 202 207 208 210

A P FOLD1

9.1 9.2 9.3 9.4 9.5

Index

Introduction 210 21 1 Background Technical/engineering basics: effect of fibres on matrix properties 218 Applications: proven and potential 226 Summary 253 References 254 257

Preface

Short fibre filled polymer composites form a relatively new family of materials, yet they are already well established in many applications. There is a vast range of materials in this category, some offering unique properties, some simply competing with other materials because of their relatively low cost. Their potential advantages are far from being fully realized and we anticipate continued growth in their use for many years to come. Research into these materials is crucial to their development and exploitation and will be for many years to come. Research continues into the design of short fibre reinforced composites and into the fundamental mechanisms that govern their behaviour, and also into methods of fabrication that will not only produce the required shape but will also result in the optimal properties being achieved. The book reflects this and is offered as an introduction to anyone starting out in research into short fibre-polymer composites. It also provides the background and bibliography for further reading. It is intended to be of value to industrial technologists who are working with these materials and are seeking further insight into their manufacture and behaviour. It is also meant to inspire materials users to consider new applications for these composites - even perhaps to formulate new ones with different combinations of properties. A special feature of the book is that it includes significant discussion on rubber-matrix fibre composites, an important sub-class of short fibre reinforced composites that is often neglected in reviews of polymer composites. We are indebted to the contributors of the chapters for their co-operation. We are grateful to them for letting us perform some editorial interventions with the aim of adding to the coherence of the text: we accept full responsibility if we have erred in this task! We wish to acknowledge also our students and colleagues at the Indian Institute of Technology, Kharagpur and the University of Newcastle upon Tyne who are responsible for stimulating and maintaining our interest in the area of short fibre filled polymer composites. We are especially grateful to those students who have conducted research in this area under our supervision and whose work forms an important part of the chapters which we have authored.

X

Preface

This book would not have appeared without the efforts of Patricia Morrison and her colleagues at Woodhead Publishing. We are indebted to our wives (Deya and Li Tong) for their expert technical help and advice during the preparation of the manuscript as well as for their patience and understanding. Finally we are thankful to our children (Barna, Dominique, Michelle, Francine and Christopher) for their cheerful acceptance of one more task competing for our attention.

S K De J R White

Contributors

M ASHIDA, 1-108 Andoji-cho, ltami 664, Japan (formerly Faculty of Engineer-

ing, Kobe University, Rokkodai Nada, Kobe 657, Japan) D M BIGG, R G Barry Corporation, PO Box 129, Columbus, Ohio 43216, USA S K DE, Rubber Technology Centre, Indian Institute of Technology, Kharagpur 721302, India A P FOLDI, C & C Consultants, 2833 W Oakland Drive, Wilmington, Delaware 19808-2422, USA B R GUPTA, Rubber Technology Centre, Indian Institute of Technology, Kharagpur 721302, India P B JANA, Super Seals India Limited, Mathura Road, Faridabad 121003, India M C H LEE, General Motors Research Laboratories, 30500 Mount Road, Warren, Michigan, 48090-9055, USA A K MALLICK, Electrical Communication Engineering, Indian Institute of Technology, Kharagpur 721302, India G B N A N D O , Rubber Technology Centre, Indian Institute of Technology, Kharagpur 721302, India J R WHITE, Materials Division, Department of Mechanical, Materials and Manufacturing Engineering, University of Newcastle upon Tyne, Newcastle upon Tyne NE1 7RU, UK S B WI LKI NSON, Materials Division, Department of Mechanical, Materials and Manufacturing Engineering, University of Newcastle upon Tyne, Newcastle upon Tyne NE1 7RU, UK

Survey of short fibre-polymer composites J R WHITE AND

I. I

S K DE

Introduction

Short fibre reinforced composites are tinding ever-increasing applications in engineering and in consumer goods. They can offer a unique combination of properties or may be used simply because they are more economical than competing materials. The matrix material is usually polymeric and in many applications they compete with unreinforced polymers. Fibre reinforcement improves the stiffness and the strength, and for many polymers it improves the toughness, though the toughness may decrease in polymers that are already tough in unreinforced form. The dimensional stability is improved and, in the case of rubbery composites, better green strength is obtained. Benefits such as creep resistance and better ageing and weathering properties may be crucial in some applications. Conductive fibres may be added to change the electrical properties. The fibre reinforced composites with the best mechanical properties are those with continuous fibre reinforcement. Such materials cannot be adapted easily to mass production and are generally confined to products in which the property benefits outweigh the cost penalty. Short fibre reinforced composites can be processed in a manner similar to the matrix. In the case of thermoplastics this means that methods such as injection moulding are available, allowing mass production of components with quite intricate shapes. Reinforced thermosets suitable for injection moulding have also been developed, whereas thermoforming (using ‘sheet moulding compound’) is another alternative for short fibre-thermoset composites. Reinforced natural and synthetic rubbers can be processed by the usual rubber processing methods such as calendering, extrusion and injection moulding, and uneconomic methods such as dipping, wrapping and laying are not required. Many of the processes cause fibre alignment, which is often beneficial. The properties are partly determined by the composition (fibre volume fraction) and partly by the processing, giving a wide range of property (and cost) combinations to which both designer and fabricator should be alert. Part of the property advantage of continuous fibre composites derives from the

2

Short fibre-polymer composites

continuous nature of the reinforcement but part is a consequence of the highly parallel fibre orientation. In short fibre composites the fibre orientation distribution is far less perfect and is often random. As a result, the degree of anisotropy is generally less than in continuous fibre composites, but it is often significant and must not be overlooked by product designers. Processing windows may have to be quite narrowly defined when making components for critical applications in order that the correct fibre orientation distribution is maintained. By adding suitable fibres and by controlling factors such as the aspect ratio, the dispersion and orientation of fibres, and the fibre-matrix adhesion, significant improvements in property can be achieved with thermoplastic, thermosetting and rubbery polymers. This book presents an up-to-date review of short fibre reinforced composites which covers thermoplastics, thermosets and rubbers. The emphasis is on the research that has underpinned their development and discussion is centred on understanding the origins of their properties. The aim is to provide an introduction to these ‘designer materials’ that will be helpful to both the component designer and the fabricator. I .2

Composition 1.2. I

fibres

The most common fibre reinforcement is glass, usually E-glass (for ‘electrical grade’). Glass fibres are usually 5-20pm in diameter and are round and fairly smooth, though their surfaces are never completely defect-free. For some short fibre reinforced composites they are provided pre-chopped in lengths of a few millimetres. For other applications they are provided already dispersed in a polymer matrix in the form of dough moulding compound (DMC) or granules. In one of the specialized forms, the granules are formed by chopping a continuous fibre reinforced pultruded rod into lengths of lOmm or so in which the fibre length equals the granule length. Yet another form of starting material is sheet moulding compound (SMC) in which short glass fibres are compacted into a continuous sheet and impregnated with a thermosetting resin at room temperature. When the compound is processed the fibres become further shortened, especially in the case of injection moulding in which severe stresses are applied to the melt. A coating (‘size’) is commonly applied to glass fibres to improve processability and to reduce damage during handling. Surface coatings may also be applied to improve the adhesion between the matrix and the fibre. It is possible to improve the properties of composites by using higher performance fibres such as carbon, boron or polyaramid fibres (KevlarTM).Such fibres are often utilized in special applications of continuous fibre composites, in which the property benefits of the high performance fibres are exploited fully. In short fibre composites, the matrix dilutes the properties much more than in continuous fibre composites and the cost of a high performance reinforcing fibre

Survey of short fibre-polymer composites

3

is rarely justified. At the other end of the cost and property spectrum, sometimes a fibrous filler that is obtained quite cheaply may be used to obtain a useful improvement in property compared with the unfilled polymer, even though the fibres are not easy to manipulate for optimum reinforcement. Such is the case with sepiolite, a clay-like mineral.'-3 Metal fibres can be used to provide electrical conductivity. The major application for thermoplastics made conductive by adding metal fibres is as housings for computers and other electrical goods requiring protection against electromagnetic interference and electrostatic discharge. Stainless steel fibres are used and loadings of the order of 1YOby volume (depending on the application) are common. Commercial grades are available based on several polymers including acrylonitrile butadiene styrene (ABS), various nylons, polycarbonate and polypropylene. This subject is discussed in Chapter 6. In the case of soft rubbery composites. cellulose fibres have been found to give better reinforcement than glass or carbon fibre^.^ The reason for this is probably that the flexibility of cellulose fibres results in less breakage during processing than happens with the brittle glass or carbon fibres which have less resistance to bending. The mixing processes applied to rubber compounds are quite intensive and are conducted while the compound possesses high viscosity, giving high bending stresses and causing severe fibre breakage with brittle fibres. The fibre length used in the preparation of rubber composites is critical. It should not be too long or the fibres will get entangled, causing problems with dispersion; if it is too short the stress transfer area is too small and the fibres do not provide effective reinforcement. Short jute and silk fibres can be used to reinforce Electrical conductivity can again be obtained by the use of additives and this is dealt with in Chapter 7. 1.2.2 Polymers

Short fibre-polymer composites based on thermoplastics, thermosets and rubbers are available and are discussed in the following chapters. The most common thermoplastic short fibre reinforced composites are based on polypropylene and nylon. Higher performance thermoplastic composites use poly(ether sulphone) or poly(ether ether ketone) (PEEK), giving higher stiffnesses and higher working temperatures. Also with thermoplastic behaviour are the thermotropic liquid crystal polymers in which the self-reinforcing property is further enhanced by the inclusion of fibres. Polyimides, which can show either thermoset or thermoplastic characteristics, are also available in reinforced form. Thermoset matrices include polyesters, epoxies and phenolics. Thermosets are generally less suitable than thermoplastics for mass production but the improved properties achieved by adding fibre reinforcement has provided added incentive to develop a suitable means of fabrication. Both natural and synthetic rubbers are candidates for short fibre reinforcement. Considerable attention is paid to short fibre-rubber composites in this book, a feature that distinguishes it from many other books in this area.

4

Short fibrepolymer composites

1.2.3

Other ingredients

In addition to the major ingredients (fibre plus polymer matrix) the composites contain several other components, the precise nature and concentration of which will normally be a commercial secret. The matrix may contain other fillers, such as talc or fly-ash, to improve the mechanical properties, and many other additives may be present as processing aids, plasticizers, stabilizers, curing agents and mould-release agents. Furthermore, the matrix may contain additives to enhance the fibre-matrix bond. The purpose and effect of such additives are dealt with as appropriate in the following chapters. 1.2.4

Preparation of short fibre composites

1.2.4. I Thermoplastics Short fibre reinforced thermoplastics are most commonly supplied in the form of granules suitable for use in injection moulding machines. The granules are roughly cylindrical, measuring 3-5 mm long and about 3 mm in diameter and the fibres are dispersed fairly uniformly within them. To make the granules, chopped fibres and unfilled thermoplastic polymer powder or granules are fed into an extruder. Mixing occurs on the passage through the extruder and the melt is passed through a die to make a lace that is chopped into granules of the required length. Fibre breakage is a significant problem during this process and the design of the screw is of some importance, deep flights being required. Twin screw extruders are often preferred and the fibres are sometimes fed into the barrel part way along so that they enter directly into melt: this is found to reduce fibre breakage. In most of the common short fibre thermoplastic moulding materials the original fibre length is around 3 mm but few fibres survive mixing and moulding intact, with most fibres present in mouldings measuring fractions of a millimetre. Another class of compound is available known as ‘long fibre’ thermoplastic, in which the granules measure about 10 mm and contain parallel continuous fibres that span the full length. These granules are produced by chopping up pultruded rod in which a continuous tow of fibres is wetted by the polymer when forced through a heated die.8 The long fibre compounds are found to behave surprisingly well in conventional injection moulding machinery and the fibre lengths are preserved more completely than is the case with compounds in which the fibres are shorter to begin with. Other methods of mixing fibres into thermoplastic polymers are given in the book by F ~ l k e s including ,~ the use of kneaders and in situ polymerization. More details of the processing of reinforced thermoplastic materials are given in Chapter 2. 1.2.4.2

Thermosets

Chopped fibres can be combined with thermosetting resin at the fabrication stage by spraying both ingredients on to a mould simultaneously. In other processes

5

Survey of short fibre-polymer composites

the fibres may be formed into a mat prior to impregnation with polymer resin. In SMC the fibres are pressed into asheet and impregnated with polyester resin then stored until the final fabrication stage when the sheet is shaped in a hot press causing the resin to cure. D M C also contains both uncured resin and filler, and has sufficiently good flow properties to be introduced into an injection moulding machine, though this is usually assisted by a 'stuffer' rather than by gravity feed as is the case with thermoplastic granules. More details of the processing of reinforced thermoietting polymers are given in Chapter 3. 1.2.4.3 Rubbers Short fibres can be mixed into rubbers using any of the rubber compounding machinery, including mills, calenders and internal mixers. Fibre breakage is a problem and the mixing conditions selected must take this, as well as other factors, into account. Mill mixing and calendering cause fibre orientation, giving rise to anisotropy and other property changes that are especially important if the final product is made this way. Details of reinforced rubber processing for various products are given in Chapters 4, 5, 7 and 8.

I .3 Morphology of short fibre reinforced polymers The detailed morphological characteristics of the different classes of short fibre reinforced polymers are described in the relevant chapters. It is necessary only to give a brief general introduction here. The properties of short fibre-polymer composites are strongly dependent on the fibre volume fraction and on the fibre orientation distribution. The fibre volume fraction is usually fairly tightly controlled, though some segregation of fibres and polymer may occur during fabrication. Mouldings sometimes have a stratified morphology in which the most prominent feature is the fibre orientation distribution which varies strongly from one depth to another within thermoplastic injection moulded short fibre-polymer composites.".' The fibre volume fraction may also vary with depth, though there are few reports of measurements of this in the literature. In addition to variations with depth, there may also be localized regions at different sites along the flow path with significantly higher or lower fibre content than the average value. Segregation of fibres in thermoplastic matrices is discussed in section 2.3.2.3 The fibre orientation distribution changes when the moulding conditions change, but it is difficult to control. Enormous benefits would be yielded if methods could be developed for exercising tight control over fibre orientation distribution in mouldings made from short fibre-polymer composites. In order to devise methods of controlling fibre orientation distribution it is necessary first to have a model for the dependence of fibre orientation distribution on processing conditions. Several studies have been made of this t o p i ~ . ' ~ - ' ' The approach by Bay and Tucker is set out in some detail' and involves a finite

'

6

Short fibre-polymer composites

difference method in which temperature and velocity variations are calculated along the flow direction and through the thickness of the part. The limitations of the method are discussed16 and the predictions are then compared with experimental results in a companion paper.' The general behaviour is found to follow the predictions reasonably well but the detailed agreement is not particularly good. The motivation for studies of this kind is to enable computer modelling of processes such as injection moulding so that the design of the moulding machinery can be optimized before it is manufactured. Thus in the case of an injection moulding tool, the geometry of runner systems, the size and positioning of gates, and the arrangement of cooling channels are chosen with the assistance of computer modelling. Even though careful experimental analysis shows that the results obtained with commercial predictive software are not particularly accurate and that further refinement should be sought, the widespread adoption of computer-aided design of moulds has resulted in a considerable reduction in the fraction of tools that require adjustment before entering service after a trial period. Most attention has been focused on the fibre orientation distribution in injection mouldings but Advani and Tucker have developed a numerical simulation of the compression moulding of sheet moulding compound and have demonstrated a reasonably good agreement between their theoretical predictions and measurements made on real mouldings.'* Many products benefit from having enhanced preferential fibre orientation in the direction of flow. The SCORIMN process has been developed to maximize this type of orientation. This is achieved by repeatedly passing melt through the mould cavity using a 'live feed' push-pull arrangement with counteracting plungers which ensures that solidification is always in a high shear zone in which orientation is During processing and subsequent fabrication of short fibre filled rubber composites, the fibres orient preferentially in one direction and the ultimate properties of the composites depend mainly on the angle between the fibre axis and the applied stress. The effects ofmill parameters such as the number of passes, nip gap and roll speed ratio have been studied and the nip gap is known to have significant effect on the fibre orientation. The greatest amount of fibre orientation takes place in the first pass. However, the orientation is also strongly influenced by the manner in which the sheet is folded and care should be taken to ensure that maximum orientation is obtained along the machine direction. Although the properties of fibre-reinforced composites depend mainly on the fibre orientation distribution, the adhesion between the fibre and the matrix is important. Loads are not applied directly to the fibres but to the matrix. For high performance, loads must be transferred effectively to the fibres, which demands a strong fibre-matrix bond. This can be controlled by surface treatments applied to the fibre and/or by modifying the polymer to provide functional groups that bond to the fibre or to the surface coating. In rubbers, the dry bonding system known as HRH, consisting of hydrated silica, resorcinol (or a resorcinol derivative) and hexamethylene tetramine (a methylene donor) is used to promote adhesion

Survey of short fibrepolymer composites

7

between fibre and rubber matrix.2 For applications requiring toughness rather than strength a weaker bond between fibre and matrix may be preferred since this will encourage the fibres to pull out of the matrix against frictional forces, giving large energy dissipation, rather than breaking at low composite deformation. Some details of the chemical nature of the bonds are given in Chapters 2 , 3 , 4 and 6. In the case of crystallizing polymer matrices, the fibre may have a strong influence on the crystal morphology. In some fibre-polymer pairings this is because crystal nucleation occurs on the fibre surface, leading to row nucleation that develops into a quite different morphology from that which forms in the absence of fibres. in other fibre-polymer pairings crystal orientation follows molecular orientation that is influenced by the fibres.22This morphology is likely to enhance the fibre-matrix bond and hence influence the properties of the composite. The changes in crystallinity and crystal orientation distribution that occur when fibres are included alter the properties of the matrix and this also modifies the composite properties.

I .4 Mechanics of short fibre reinforced polymers In a fibre-polymer composite the fibres are stiffer than the matrix and the proportion of the load that they support is greater than their volume fraction. The overall elastic properties of a composite are relatively easy to compute from the elastic properties of the components when the fibres are continuous and parallelz3and it is worth considering briefly this special case before moving on to the much more challenging problem presented by short fibre-polymer composites. The derivations of the relationships presented below are found in any standard text on composites (e.g. the book by P i g g ~ t t ~ and ~ ) will not be reproduced here. In the following discussion the properties of the fibre will be indicated by subscript f, the polymer matrix by m and the composite by c. Directions in the composite are indicated by subscript notation as described, for example, by Young and L0ve11.~~ We have chosen to represent the fibre axis direction by ‘3’: the reader is warned that this convention is not universally applied and some authors, including P i g g ~ t t , ’use ~ ‘1’ to represent the fibre axis direction. Thus if the Young’s modulus is represented by E and the volume fraction by V the Young’s modulus in the direction of the fibres, Ec,33,is:

This expression follows from the assumption that uniform strain is present in the composite when it is loaded in this way. When an aligned fibre composite is loaded in the transverse direction, most of the deformation is taken up by the soft polymer phase and it is assumed that a state of uniform stress prevails. This leads to the following expression from which can be calculated the transverse Young’s modulus for the composite ( E c , l

8

Short fibre-polymer composites

If typical values for E,, Em, V, and V,, are substituted into equations [1.1] and [1.2], it is confirmed that the aligned fibre composite is very anisotropic. Once the values for Ec.33 and E c , , are known, the Young's modulus, 4 , at an angle 0 to.the fibre axis, can be calculated and is given by:

,

1

cos40

_ -- -+ E, E c . 3 3 where the shear modulus G

and

\'13

(Gf,

c.:)

sin20 cos20

+ sin4 0 ~

~1.31

EC.1 I

is given by:

is Poisson's ratio defined in the following way: \ l l 3

=

-

strain along libre axis transverse strain

[I 1.51

L 3is likely to be small because deformation is much easier transverse to the fibre axis than along it. Strengths are very sensitive to the presence of flaws and therefore to the manufacturing process, and are more difficult to predict than stiffnesses. The analysis of the mechanics of short fibre composites is much more difficult than for continuous aligned fibre composites. There are two reasons for this. Firstly the stress transfer between fibre and matrix is not uniform along the fibre, and there are end effects that can be neglected in continuous fibre composites but that are important in short fibre composites. Secondly in short fibre composites the fibres are never exactly parallel and may even have random orientation. Methods to take account of the fibre orientation distribution have been proposed: they demand considerable computation and require that the fibre orientation distribution is measured. These problems are considered in turn in the following sections.

1

1.4.I

Stress transfer at the interface between a short fibre and the matrix

Consider a single iso!ated short fibre of diameter rl within a continuous matrix. If a homogeneous strain is applied to the matrix parallel to the fibre axis, load is transferred to the fibre by means of shear stresses at the interface (Fig. 1.1). The shear stress varies along the fibre axis and, a s a consequence, the axial stress within the fibre varies. If a small tensile deformation is applied parallel to the fibre axis (s-axis),the stress in the fibre at s. o,(.s), and the interfacial shear stress acting on the surface of the fibre, x i s ) , can be related by solving the force balance equation, which reduces to:23

-

9

Survey of short fibre-polymer composites

TJX)

\

1

I

I

I

I

I

I

I

I

a,(xl-*

I

-a,+

I

I

I

I

I

I

dacdx dx

I

I I

x+dx

X

1.1 Schematic representation of a fibre embedded in a polymer matrix. When strain is applied to the composite, stress is transmitted to the fibre via the interfacial shear stress, T ~ .

t

RI cc$;____)

........................ -----

----------+------------------

-

........................ 7

1.2 Schematic for Cox’s shear lag analysis. The fibre, diameter d, is embedded in a polymer matrix, the influence of which is considered to extend as far as the surface of a cylinder, radius R, where 2R is the average fibre separation.

The stress distribution along the fibre is derived using Cox’s ‘shear-lag a n a l y ~ i s ”in ~ which the fibre is considered to be surrounded by a cylinder of polymer matrix of radius R where 2R is the average fibre separation distance (Fig. 1.2). If the shear stress on a cylindrical surface radius r ( < R ) in the matrix is T then for a fibre of length I a n d diameter d, force balance requires that:

2 n d ~= ndlri i.e. T

dT 2r

=-

c1.71

If u is the axial displacement at axial position x on a cylindrical surface within the matrix at radius r, then the shear strain in the cylinder of matrix is given by:

10

Short fibrepolymer composites

au ar

T.

d Ti 2r G,

-- --_ _

G,

At the fibre surface u = u, and r = d/2 while a t r = R, u measured displacement in the composite). Thus we have:

L1.81 = ti,

(the overall

i.e.

[l.lOj

Differentiating with respect to x gives: [1.11]

where E , and E , are the strains in the composite and the fibre respectively. Differentiating equation [ 1.63 with respect to x gives: d20, dx2

~

4 dri d dx

[1.12]

F r o m equations C1.111 a n d C1.123 we have: [1.13]

where c, has been replaced by o,/E,. T h e general solution to equation C1.133 is of = E,E,

+ C sinh(Ax) + D cosh(Ax)

C1.141

where:

d

E,ln(2R/d)

[l.lS]

If we now place the centre of the fibre a t x = 0 and assume that no axial stress is transmitted across the ends of the fibre then o f = 0 at x = - and x = 3,a n d the solution becomes: [1.16]

T h e stress in the fibre increases from the fibre end and reaches a maximum a t the centre unless failure occurs. T h e interfacial shear stress is obtained by differentiating equation C1.161 (see equation C1.61):

Survey of short fibrepolymer composites t. =

’

AdE,E, sinh Ax 4cosh( A;)

II

C1.171

Note that equations [1.161and [ 1.171show the dependence of the stress transfer within the composite on the fibre aspect ratio, l/d (recall that A is a function of d ) and confirm that it should be as high as possible. The maximum magnitude oft, (equation [ 1.171)occurs at the fibre ends and, as a consequence, this is where the first event leading to failure will occur by debonding from the matrix or by shear failure in the matrix. Thus ifthe composite is deformed sufficiently for this to occur, a slipped region develops and equations [l.l6]and C1.171 nolongerapply. Adjacentto theslipped region thefibrestressis obtained by solving equation [1.6] for constant t i ( = T?, say). The result can be rearranged to give the length of the slipped region at each end of the fibre, I,, in terms of the stress in the fibre in the unslipped region (taken to be a constant value, c,,J i.e. I, = o ~ . ~ ~ / ~The T Tfibre . will break if ( T ~ reaches , ~ the breaking stress, but this can happen only if the fibre length exceeds a critical length, 1cr,,. = 2 1 s,cr,,. = cr,crild/2~:.Generally, for the best mechanical properties, the fibre length should exceed I,,,,, but if the composite is overloaded fibre fracture will occur until the fibres degrade to this value. The analysis presented above shows that both the fibre length and the aspect ratio should be as large as possible. The dimensions of the short fibres for reinforced composites must be chosen accordingly and steps must be taken during processing to maintain lengths at adequate levels. 1.4.2 Young’s modulus of a short fibre-polymer composite

Mathematical formulae for predicting the stiffness of a composite possessing an idealized fibre orientation distribution are well established. The most commonly studied distributions are (i) all fibres parallel to the test direction, (ii) all fibres transverse to the test direction and (iii) random orientations. The models also assume that the matrix is fully isotropic. Such models are therefore not applicable to injection moulded composites. Several mathematical models have been developed to predict Young’s modulus from the fibre orientation distribution; the one presented below2‘ has been developed from methods described previously by other workers. I t is essentially an adaptation of the aggregate theory of Brody and Ward,27 incorporating the depth dependency of both the fibre orientation and the matrix stiffness. In injection moulded bars the overall anisotropy is expected to result from the preferred orientation of fibres and from anisotropy within the polymer phase. Here we attempt to account for the latter by relating the matrix anisotropy to the local fibre orientation. In the aggregate theory the composite is considered to behave as an aggregate of identically anisotropic sub-units, each with the elastic properties of fully oriented material. The matrix is assumed to be anisotropic with transverse

12

Short fibrepolymer composites

I'

1.3 The relationship between the global reference frame of the composite bar and the local sub-unit reference frame. MFD = mould flow direction.

isotropy, whereas the glass fibres are fully isotropic. The sub-units are therefore transversely isotropic and an orientation distribution D(+)is considered adequate to describe the distribution of sub-unit orientations. The relationship between the global reference frame of the composite (as defined by the injection moulding process) and the local sub-unit reference frame (as defined by the fibre alignment) is shown in Fig. 1.3. The properties of the sub-unit are represented in the usual mannerz8 using unprimed suffixes. The elastic constants of the composite ('aggregate') are represented using primed suffixes. It has been shownz8 that the aggregate axial compliance '53,38is given by: s3,3,

= IIS, 1

+ fzS,, + 13(2'513 + '544)

E1.181

where: &+)sin4 4 d 4 1, = < sin4+ > =

11.193

Survey of short fibre-polymer composites

lo*''

13

D(4)cos4+d+

1, = < C O S " ~> =

[;"

[1.203

D(4)d+

I, = < sin2t$ cos2 4 >

C1.21)

D ( 4 ) takes the form: [1.221

where N ( 4 , A+) is the number of fibres lying at orientations between (+ - A4/2) and (4 + A4/2). In the calculations performed by O'Donnell and White26 the values taken were A+ = 10" and 4 = 5", 15", 25"... 85". I , , I , and 1, must be calculated by numerical integration over the measured fibre orientation distribution. The compliances of the sub-unit were calculated from the following expressions:28

s,, =

1 -'

E,'

1 LT s,, = -.E,1' 's44 -; s,, -- GL ET

where ET is the transverse modulus of the sub-unit, EL is the longitudinal modulus of the subunit, G, is the longitudinal shear modulus of the sub-unit and vLT is the Poisson's ratio of the sub-unit. The elastic constants of the sub-unit were determined from the Halpin-Tsai

equation^:'^ [1.231 [1.241

where

The elastic constant S,, was calculated using:29

G

+ GL

G: L -

[1.251

2

where G,=G,+

I

-

1

"F

1-

v,

[1.261

14

Short fibrepolymer composites

[ 1.271

Within a short fibre reinforced injection moulding there exists a distribution in fibre length. This must be incorporated into the aggregate theory and this can be done by taking the maximum and minimum values for fibre length, I,, to give upper and lower bounds for the composite stiffness.z6 O’Donnell and Whitez6 showed that the calculations of composite stiffness are not very sensitive to the fibre length when it is in the range found in the typical injection moulded composites used in their studies. Some of their results are given in Chapter 2. O’Donnell and White perfomed a layer-wise analysis and obtained Young’s modulus as a function of depth through the mouldings. The value of the Young’s modulus of the matrix inserted into the equations was taken to be that measured in mouldings made from the same polymer in unfilled form. This varied with depth and it was shown that this produced better agreement with the measured Young’s modulus d’stribution than when the calculations used a uniform matrix Young’s I .5

Measurement of fibre orientation distribution

It is evident f i J m the discussion presented in section 1.4 that the characteristic that is most important in determining composite mechanical properties is the fibre orientation distribution. In some cases it is sufficient to have a qualitative assessment of the fibre orientation distribution, as, for example, near a weld line in a multiply-gated injection moulding. In mouldings having relatively simple shapes and gating arrangements it is an advantage to have quantitative measurement methods of fibre orientation distribution so that an objective assessment may be made of the effect of different processing procedures designed to control the fibre orientation distribution. It is equally important to be able to determine the parameters describing the orientation distribution for insertion into the equations that describe the mechanical behaviour. This is normally done by image analysis, which can also provide measurement of the fibre length distribution, also needed for the mechanics analyses. I .5.I

Image analysis of seaions

The first step is normally to prepare a section of the sample cut in a chosen direction, often normal to the flow direction in an injection moulding (but see discussion later on). The section is then polished using metallographic procedures, developing improved polishes through the successive use of abrasive papers followed by a polishing wheel with cloth-borne abrasive particles of diminishing s i ~ e . ~ ’ .This ~ ’ requires much care with unfilled polymers because of their softness and tendency to heat when abraded; the use of a cooling mixture of

Survey of short fibre-polymer composites

15

water and detergent is re~ommended.~' With fibre reinforced materials the large difference in stiffness or hardness between the fibre and matrix can lead to problems and care must be taken to ensure that differential removal of the two phases does not occur and that fibre breakage does not take place. Toll and Andersson claim that uneven cutting occurs if the final polish is made with particles that are too fine and recommend a final polish with 2pm diamond particles.32 Examples of sections through short fibre reinforced composites are given in Chapter 2 (thermoplastics) section 2.3.2.3, Fig. 2.2, and in Chapter 3 (thermosets) section 3.6, Fig. 3.2 and 3.3. When viewed in the light microscope, the contrast between the fibre and the matrix may not be sufficient, especially for automated image analysis, and when this is thecase an etch may be used, generally to erode the polymer matrix slightly in the case of thermoplastics. Alternative procedures for thermoset matrix composites include etching the sample either with hydrofluoric acid (which attacks the fibres) or with oxygen plasma. Another method for enhancing contrast is by applying a vapour deposited coating, usually gold or platinum instead of an etch. Touching fibres can present a problem in automatic image analysis but instructions can be included in the software for recognition and suitable a c t i ~ n . ~ ' ? ~ ~ The most common method of fibre orientation measurement is by using the light microscope and an image analysis package. The image analyser measures the aspect ratio (maximum/minimum diameter) of the elliptical image of each fibre section in the field of view; the reciprocal of this quantity equals cos 4, where 4 is the angle between the fibre axis and the normal to the sampling plane. A field typically contains 512 x 512 pixels on a video screen. Each pixel is examined separately by the image analyser and described by its x and y co-ordinates plus a digitized grey level which describes its shade usually in 256 steps between black (0)and white (255). The grey level assigned to an individual pixel may be altered to enhance the image contrast. The enhanced grey image is then available for image analysis. The orientations of the fibres are determined from their shape and the orientation of their elliptical intersections on the plane of the section. The error involved in digitizing the diameters when they are expressed in terms of pixels is important when the fibre is closely parallel to the normal to the because of the shape of the cos 4 function for small values of 4, and for typical image magnifications the error in 4 is of the order of 10". To avoid this problem, Hine et al recommended the use of oblique sections, cut at an angle between 35" and 75" to the flow direction, when a strong preferred orientation is present.36 The objective of the image analysis is usually to produce a volume average of the fibre orientation distribution rather than a description of the fibre orientation distribution of those fibres that happen to intersect the sampling plane. Those fibres that are oriented normal to the sampling plane are most likely to intersect it and a weighting function of l/cos$ or similar should be applied to compensate for this.32*34*37 Fibres that are inclined to the normal to the sampling plane and are intersected near to one end appear as truncated ellipses, and ways of dealing

16

Short fibre-polymer composites

with this are discussed by Bay and Tucker34 and by Fischer and E ~ e r e r . ~ * Another problem occurs when a fibre that lies nearly parallel to the sampling plane is curved along its length and this is discussed by Toll and A n d e r ~ s o n . ~ ~ Recent developments in image analysis coupled with the rapid measurement of thousands of elliptical intersections, render sampling errors negligible, though this is not generally available on commercial image analysis package^.^^.^^ In another development, the use of confocal microscopy, producing images from different selected planes beneath the polished surface, allows 3D orientations to be measured from a single scan;35 again this procedure is not in routine use. Conventional commercial image analysis equipment has been used by several workers to investigate the orientation of fibres that intersect a plane in a short fibre The fibre orientation can be represented by an orientation parameter as recommended by Fakirov and F a k i r ~ v a , ~who ’ used two orientation parameters introduced by Hermans4’ to describe the orientation in crystalline polymers. Assuming that there is a planar distribution of fibres, the Hermans orientation parameter f, is of the form: fp=2 - 1 where:

The modified Hermans orientation parameterfb is of the form: fb=2cos2<4> -1 where:

+ is shown in Fig. 1.3. The angle that the major axis of the ellipse makes with the y direction is equal to 0 . The angle that the fibre axis, OF, makes with the major flow direction (the x-axis) is equal to and is determined by calculating aspect ratios of the fibre intersections:

+

cos+

r R

=-=

length of minor axis length of major axis

The orientation parameters are scaled so that for a composite material where the fibres are randomly oriented f, andf, are both zero. If the fibres are aligned parallel to the flow direction then bothf, andfp = 1 and if the fibres are aligned perpendicular to the flow direction.f, = = - 1. For a section taken at an angle TX to the vertical axis (as in the work by Hine et it is possible to obtain both 8 and using simple trigonometric transformations.

,rp

+

17

Survey of short fibrepolymer composites

1.5.2 Contact microradiography

If a thin section can be made, the fibre orientation distribution can be revealed by microradiography whereby X-rays are passed through it and their transmission is recorded on an X-ray sensitive film.43The shadows thrown by the fibres d o not permit the assessment of their inclination to the plane of the section. As an alternative to performing image analysis on the microradiographs, McGee and McCullough suggested using optical diffraction to obtain measurements of average orientation^.^^

I .6

Properties of fibre reinforced polymers I .6.I

Average properties

Table 1.1 lists some of the key properties of structural materials. Short fibre reinforced thermosetting polymers are represented by SMC and DMC, and short fibre reinforced thermoplastics by GFPP and GFN6,6. Much of the property range given for SMC and DMC is due to differences in composition from different sources (largely undeclared). The same is true with thermoplastics: glass fibre reinforced nylon 6,6 with 20% by weight of glass is much closer in property to the polypropylenecompound with a similar glass content than it is to GFN6,6 with 40% by weight of glass. The stiffness can be almost doubled by replacing glass fibres by carbon fibres; this also causes significant improvement in the strength and the density (which is lower in carbon fibre reinforced compounds). The true value of reinforcement can be judged first of all by comparing the composites with the corresponding unreinforced polymers in Table 1.1, rememTable 1.I . Properties of selected materials Material Aluminium Steel E-glass Carbon fibre Kevlar-49 Nylon 6,6 Polyesters Polypropylene Polystyrene Rubber SMC, D M C G F P P (20%) GFN6,6 (40%) Reinforced rubber

Young’s modulus, G N m-*

70 210 70-76 230-400 130 2- 5 1-5 0.9-3 3 0.001 8-1 5 4 14.7 0.14.4 ~

Strength, M N m - ’

Density, kgm-’

50-70 4W3000 1500-3000 2100-32OO 2700-3000 70 45-90 3MO 40

2700 7800 2500 1750-2200 1450 1140 IlW1400 890-910 1050 850-900 1700-2600 1040 1460 1050-1500

60- 140 65 220 20-50 ~

~

~~

~~

~

SMC - sheet moulding compound; D M C - dough moulding compound; G F P P - glass fibre reinforced polypropylene (20% glass by weight); GFN6,6 - glass fibre reinforced nylon 6,6 (40% glass by weight); Reinforced rubber - typical values for loadings of 1WO% by volume.

18

Short fibre-polymer composites

bering that the reinforced materials can be processed using the same methods. Secondly, when assessing the suitability of the materials for certain structural applications it is often instructive to divide the key property by the density to obtain a figure of merit or ‘performance index’. For example, aerospace (and, to a lesser extent, automotive) applications normally require stiffness and at the same time weight saving. When this is taken into account the polymer composites can often compete very strongly with high performance metal alloys. The derivation of appropriate performance indices and their application is discussed by Ashby.*’ 1.6.2

Depth variation of properties

The properties discussed above are those averaged through the section of the moulding. If there is fibre segregation or if the fibre Orientation distribution varies through the depth of the moulding then there will be a variation in property through the thickness. For example, Young’s modulus and thermal expansion coefficient both depend sensitively on the fibre orientation distribution and will vary through the depth accordingly. A straightforward average can be expected for the overall values in the preferred fibre orientation direction but other properties may not be as easily derived. The stiffness in bending is more sensitive to the distribution in Young’s modulus. It is given by the summation through the depth of the product of the second moment of area and the (depth-varying) Young’s modulus, i.e. bjE(z)zdz where b is the width of the bar. This is the key parameter when designing sandwich structures for load-bearing applications, for which the departure of its value from that calculated assuming a uniform average Young’s modulus is very great indeed. In a similar way, a reinforced moulding with a high Young’s modulus near the surface is stiffer in bending than would be predicted if only the average Young’s modulus were available. Sharp changes in property occur whenever there is a sudden change in the fibre orientation distribution. Thus in injection mouldings with a stratified morphology there will occur narrow boundaries between different regions with a significant mismatch of properties such as Young’s modulus and linear thermal expansion coefficient. Under certain loadings or thermal conditions this may promote delamination and such a morphology should be avoided if possible. Examples of measurements of depth variation in properties are given in Chapters 2 and 3.

I .7 Conclusions Reinforcement of polymers using short fibres can lead to unique property combinations. Even if the properties of the short fibre-polymer composite are not unique, reinforcement can produce significant property benefits at a modest cost. In most cases the composites are processed easily using conventional polymer or rubber processing methods. Further improvements in property could be achieved by exercising more control over the fibre orientation distribution. Variations in fibre orientation distribution commonly occur within mouldings,

Survey of short fibre-polymer composites

19

particularly through the wall thickness. This is unlikely to produce optimal properties and may even lead to failure. References 1 Acosta J L, Morales E, Ojeda M C and Linares A, J Muter Sci, 21 (1986) 725. 2 Acosta J L, Morales E, Ojeda M C and Linares A, Angew Makromol Chem, 138 (1986) 103. 3 Morales E and White J R, J Muter Sci, 23 (1988) 4525. 4 Murthy V M and De S K, Polym Eng Revs 4 (1984) 313. 5 Murthy V M and De S K , J Appl Polym Sci 29 (1984) 1355. 6 Akhtar S, De P P and De S K, J Appl Polym Sci 32 (1986) 5132. 7 Setua D K and De S K, Rubber Chem Tech 56 (1983) 804. 8 Crosby J M, Chapter 5 in Thermoplastic Composite Materials, Ed L A Carlsson, Elsevier, Amsterdam (1991) 139. 9 Folkes M J, Short Fibre Reinforced Thermoplastics, Research Studies Press/Wiley, Chichester (1982). 10 Darlington M W, McGinley P L and Smith G R, J Muter Sci, 11 (1976) 877. 11 Bright P F , Crowson R J and Folkes M J, J Muter Sci, 13 (1978) 2497. 12 De Frahan H H, Verleye V, Dupret F and Crochet M J, Polym Eng Sci, 32 (1992) 254. 13 MatsuokaT, Takabatake J-I, Inoue Y and Takahashi H, Polym Eng Sci, 30 (1990) 957. 14 Matsuoka T, Polypropylene: Structure, Blends and Composites, Vol3:Composites, Ed J Karger-Koksis, Chapman & Hall, London (1995) 1 13. 15 Lockett F J, Plast Rubber-Proc, 5 (1980) 85. 16 Bay R S and Tucker C L , 111, Polym Compos, 13 (1992) 317. 17 Bay R S and Tucker C L, 111, Polym Compos, 13 (1992) 332. 18 Advani S G and Tucker C L, 111, Polym Comp, 11 (1990) 164. 19 Allan P S and Bevis M J, Plast Rubber Proc Applics, 7 (1987) 3. 20 Gibson J R, Allan P S and Bevis M J, Plast Rubber Inr 16(5) (May 1991) 12. 21 Murthy V M and De S K, Rubber Chem Technol, 55 (1982) 287. 22 Campbell D and White J R , Angew Makromol Chem, 122 (1984) 61. 23 Piggott M R, Load Bearing Fibre Composites, Pergamon, Oxford (1980). 24 Young R J and Love11 P A , Introduction to Polymers, 2nd Ed, Chapman and Hall, London (1991). 25 Cox H L, Brit J Appl Phys, 3 (1952) 72. 26 O’Donnell B and White J R, Plast Rubber Compos Proc Applics, 22 (1994) 69. 27 Brody H and Ward I M, Polym Eng Sci, I 1 (1971) 139. 28 Curtis A C, Hope P S and Ward I M, Po1,ym Combos, 3 (1982) 138. 29 Halpin J C and Tsai S C, Eflects of environmental factors on composite materials, AFML-TR-67, (1969). 30 Sawyer L C and Grubb DT, Polymer Microscopy, Chapman and Hall, London, New York (1987). 31 Bartosiewicz L and Mencik Z, J Polym Sci Polyrn Phys Ed, 12 (1974) 1163. 32 Toll S and Andersson P - 0 , Composites, 22 (1991) 298. 33 Clarke A, Davidson N and Archenhold G. Proc Int Con! Transpuring ’91 vol. 1,IOS Press, California (1991) 3 1. 34 Bay R S and Tucker C L, 111, Polym Eng Sci 32 (1992) 240. 35 Clarke A, Davidson N and Archenhold G, J Microsc, 171 (1993) 69.

20 36 37 38 39 40 41 42 43 44 45

Short fibrepolymer composites Hine P, Duckett R A, Davidson N and Clarke A R, Compos Sci Technol, 47 (1993) 65. Moginger B and Eyerer P, Composites, 22 (1991) 394. Fischer G and Eyerer P, Polyrn Compos, 9 (1988) 297 . Clarke A, Davidson N and Archenhold G, Trans Roy Microsc SOC,1 (Proc Conf Micro '90, London), Adam Hilger, Bristol (1990) 305. Fakirov S and Fakirova C, Polyrn Compos, 6 (1986) 41. Vaxmann A, Narkis M, Seigmann A and Kenig S J, J Muter Sci Letts, 7 (1988) 25. Hermans P H, Contributions to the Physics of Cellulose Fibre, Elsevier, Amsterdam (1946). Darlington M W and McGinley P L, J Muter Sci, 10 (1975) 906. McGee S H and McCullough R L, J Appl Phys 55 (1984) 1394. Ashby M F, Materials Selecfion in Mechanical Design, Pergamon Press, Oxford (1992).

2 Short fibre filled thermoplastics J R WHITE

2. I

Introduction

The addition of fibres to thermoplastics produces significant improvements in mechanical properties. As with most fibre composites both the stiffness and the strength are much higher than in the corresponding unfilled polymer and the toughness is often improved as well. The resistance to creep is much higher in fibre reinforced grades and the dimensional stability is generally better, though mouldings made from fibre reinforced polymers are sometimes more prone to warping than mouldings made from unfilled polymers. The inclusion of fibres can often provide a useful increase in the maximum service temperature. These benefits are obtained without sacrificing the mouldability of the materials and they can be used in conventional injection moulding machines with tools designed for unfilled thermoplastics. Normally no attempt is made to control the positioning of fibres and their orientation distribution is not ideal, but the overall distribution is generally favourable, though not optimal. Wear of the processing machinery caused by the abrasive filler particles can be considerable and should be taken into consideration when determining the commercial aspects of material selection. Injection moulding permits rapid production of components with complex shapes and reinforced thermoplastic mouldings are used in a wide variety of applications. Examples of gears and drive shaft components are shown in Fig. 2.1.

2.2 Materials 2.2. I

Polymers

Most thermoplastics are candidates for fibre reinforcement, although the end use has a significant influence on the list of common grades that have been made available commercially. Fibre filled grades are opaque, and this excludes them from many applications. Although products are often made from polystyrene because of its high transparency, its cheapness and excellent processing properties have resulted in the development of glass fibre reinforced grades for other

22

Short fibre-polymer composites

2.1 Gear and drive shaft mouldings made from fibre reinforced thermoplastics. (Photograph courtesy of ICI Chemicals and Polymers Limited.)

quite different applications. Reinforced acrylics are not as well known. Short fibre reinforcement is used to enhance the properties of other amorphous polymers including polycarbonate, poly(ether sulphone) and poly(ethy1ene terephthalate). Unfilled crystallizing thermoplastics are normally translucent or opaque and the addition of fibres is rarely restricted by optical property requirements. The most common short fibre reinforced polymers are polypropylene and the nylons (mainly nylon 6,6 and nylon 6). For high performance, composites based on poly(ether ether ketone) (PEEK) or poly(ether ketone) (PEK) are available. To maintain good processing properties, with melt flow behaviour suitable for injection moulding, it is sometimes necessary to use a lower molecular weight than that preferred in an unfilled grade. Finally it is noted that fibre reinforced thermotropic liquid crystal poiymers have been made available commercially. The combination of the fibre reinforcement and the self-reinforcing property of the liquid crystal polymer produces a particularly stiff composite. 2.2.2 Fibres

The moulding materials are normally supplied in the form of granules with the fibres already mixed into the polymer, usually at a concentration within the range 10-40% by weight, which corresponds very roughly to 5-20% by volume, depending on the densities of the polymer and the fibres. The granules are normally cylindrical with length 3-5 mm and N 3 mm diameter, and the fibres

Short fibre filled thermoplastics

23

may span their length or, more usually, are much shorter. The fibre length becomes significantly degraded during pro~essing’-~ but the residual length still provides useful reinforcement. A class of materials has been developed in which the granules and the fibres measure 10 mm in length, and, with some relatively simple modifications to the moulding conditions and the geometry of the flow path, it is possible to preserve quite high fibre lengths in the moulded products, giving superior reinforcement. These compounds are often referred to as ‘long fibre’ composites (e.g. Verton, originally developed by ICI and now available from DuPont) but this is simply to distinguish them from the more common short fibre composites and it is entirely appropriate to include them in this text in which the term ‘short fibre’ is meant to exclude continuousfibre composites, but not to restrict consideration to fibre lengths lower than some arbitrary value (say 5 mm). The mouldings shown in Fig. 2.1 are made from Verton and it is evident that the use of relatively long fibres does not prevent the manufacture of quite intricate shapes. Glass fibres are overwhelmingly the most common fibre reinforcement, both in short fibre composites and in ‘long fibre’ composites. Their properties are good and they are available at consistent quality and reasonable cost. Carbon fibres have higher stiffness and strength but their substitution for glass fibres produces only modest improvements in the properties of the composite because the matrix properties dominate. Carbon fibre reinforcement is therefore used only in applications where the need for performance outweighs the need for cost control, such as aerospace components and specialized sports equipment. There are no special advantages offered by ceramic fibres as reinforcement for thermoplastics, though the prospect of thermoplastics containing fibres made from high temperature superconductor ceramics is intriguing. Metal fibres can be used to provide electrical conductivity’ (see Chapter 6). The use of polymer fibres is restricted by the desire to fabricate using melt processing. There are many polymers with processing temperatures low enough to use fibres made from lyotropic liquid crystal polymers such as polyaramids (e.g. poly(parapheny1ene terephthalamide), Keular, or poly(parapheny1ene benzobisthiozole), PBZT),6-8 but the cost of these materials is rarely justified except when they are used as continuous fibres, when the maximum benefit can be obtained. 2.2.3 Particulate reinforcement

The use of particulate reinforcement might appear to be outside the scope of this book, but in the case of thermoplastics it demands consideration because particulate reinforcement is often used for the same reasons as fibre reinforcement and the materials produced offer similar properties and often compete in the market-place. In some cases the particulate is fibrous, as for example, sepiolite, a clay-like mineral which has been investigated as a filler for p o l y p r ~ p y l e n e . ~ * ’ ~ Other minerals such as talc and mica come in the form of platelets and share with short fibres the ability to align during moulding as a consequence of flow, and produce a generally similar anisotropic reinforcement. Mouldings reinforced

24

Short fibrepolymer composites

with glass fibre tend to be prone to warping but this is less of a problem with polypropylenefilled with mica;' ' the inclusion of mica in combination with fibres can help to counteract warpage.12 Composites made from equiaxed particulates such as chalk generally display less reinforcement and are less anisotropic than unfilled injection mouldings, but particulate fillers are effective in improving dimensional stability and resistance to elevated temperatures. 2.2.4 Coupling agents

Coupling agents are frequently employed to improve the adhesion between the filler and the polymer. Some of the basic chemistry of coupling systems is dealt with in Chapter 3 (section 3.7.1). For some combinations of filler and polymer the composition of the coupling agent (and hence its detailed action) is protected by commercial secrecy. The most popular group of coupling agents are the silanes which can bond both to the inorganic glass or mineral filler and to the organic polymer. The filler surface usually has hydroxide groups to which the coupling agent bonds, forming a bridge across to the polymer matrix (see section 3.7.1). Acrylic acid is often used to improve fibre-matrix adhesion13 and a combination of polypropylene with acrylic acid grafted on to it with silane-treated mica has been studied by Chiang and Yang.14 The inclusion of a coupling agent normally produces a significant improvement in stiffness and strength. The improvement in flexural strength is typically from 10 to 90%." One of the main causes of weakening of fibre composites is the destruction of polymer-matrix adhesion by the ingress of water, and particular attention has been paid to developing coupling agents that resist this a t t a ~ k . ' ~ The effectiveness of a coupling agent is quite sensitive to the method used to fabricate the composite. The improvement in property is much greater with compression moulding than with injection moulding.' This is not simply because the injection moulding has better properties to begin with, but is largely because the high shcar conditions that prevail during moulding discourage the formation of the interfacial bonds. Ionomeric coupling agents have been developed that partly overcome this problem. " Coupling agents are also used to advantage with particulate fillers, including mica, talc and glass flake^.'^.'^ * O 2.3

Fabrication - injection moulding 2.3. I

Introduction

The vast majority of articles made from fibre reinforced thermoplastics are produced by injection moulding. Most injection moulding machines are of the reciprocatingscrew type in which the polymer is melted in a barrel within which is an Archimedian screw. The material in the form of granules is fed into the barrel via a hopper and falls on to one end of the screw. The flights at this end of the screw are deep to permit conveyance of solids and as the screw turns the material

Short fibre filled thermoplastics

25

is moved forwards into a zone heated by electric resistance heaters attached around the outside of the barrel. The polymer melts and further heating is obtained from mechanical work as the screw turns through the viscous melt. At the forward end, the shank of the screw has a larger diameter so that the channel defined by the flights is shallower and the work done on the material intensifies. This assists mixing but also increases fibre length degradation. As material is fed forward the screw is allowed to move backwards along the barrel axis and a charge of homogenized melt gathers at the front end of the barrel. The resistance to the motion of the screw during this part of the process, or screw-back pressure, can be adjusted and is likely to influence the degree of fibre degradation. When sufficient material is present to fill the mould cavity plus the runner system the screw is thrust forwards and, acting as a ram, propels the melt into the mould via a nozzle that is held tightly against the entrance to the mould. The injection speed, or rate at which the screw is thrust forward, is normally under the control of the operator and is another parameter that may influence fibre degradation. The material in the mould must now be allowed to cool until it is sufficiently solidified. The force on the screw is maintained for a significant fraction of the cooling time so that thermal shrinkage of the melt can be combated and the mould is kept ‘topped up’. This is often most satisfactorily achieved using a holding pressure that is less than the injection pressure. The melt is admitted into the mould cavity from the runner system through a narrow constriction called a gate (or a set of gates in the case of large mouldings or mouldings with complicated shapes). Once the material in the gate has frozen, no more topping-up can occur, the holding pressure can be released, and the screw can start to turn again and prepare the next charge. Further cooling time must be allowed to elapse before the moulding has solidified sufficiently to be ejected. The common commercial grades of fibre reinforced thermoplastics are designed for use in standard injection moulding machines and work satisfactorily at temperatures, pressures, injection rates and timings similar to those used with unfilled thermoplastics. There are differences between fibre reinforced materials and unfilled thermoplastics, however, and these must be heeded. Firstly, it should be noted that the fibres are generally abrasive and cause wear of the injection moulding machine and tool; this should be taken into account when selecting the material for a product, though the choice is usually dictated by the property requirements of the product, and the end-use properties are normally very different for unfilled and fibre filled materials. Secondly, it is prudent to attempt to minimize fibre length degradation by appropriate design of the tool and selection of operating conditions. This is particularly important when moulding ‘long’ (10mm) fibre reinforced grades, and is discussed later. 2.3.2 Morphology in injection moulded and fibre filled thermoplastics

2.3.2.I Introduction Unfilled injection moulded thermoplastics generally contain a multilayered structure. Injection mouldings made from glassy polymers usually possess a

26

Short fibre-polymer composites

skin-core morphology with a surface region of the order of 0.1-0.4 mm thick that is quite different from the interior. The reason for this is that the surface cools rapidly when the melt contacts the mould cavity wall, whereas the material in the interior cools much more slowly because of the low thermal conductivity of the polymer. The cooling rate near the surface is normally too fast to permit much recoil of the molecules that have become oriented during the mould filling process. Thus, in the skin, much of the flow-induced molecular orientation is retained in a 'frozen-in' state. As long as the moulding remains below the polymer glass transition temperature, T,, there will be very little change in the conformation of the molecule backbone, though some localized reorganization may occur, leading to density changes. In the interior, cooling takes place sufficiently slowly to allow the molecules to recoil significantly, though not necessarily completely, and the core is characterized by a much lower degree of orientation. The skin-core morphology is generally more evident in semi-crystalline polymers than in non-crystalline polymers. This is because the differences in molecular orientation through the depth of the moulding lead to different crystal morphologies, and sometimes even different crystal structures in the case of a polymorphic polymer such as polypropylene. Near the surface the pronounced molecular orientation usually causes a preferred orientation in the crystals that grow there. In the interior, however, the molecules recoil before they cool sufficiently to form crystals and when crystallisation does eventually commence the melt is effectively isotropic, and, in unfilled polymers, equiaxed spherulites form. The spherulite size may vary with cooling rate and may change through the depth of the moulding. In addition to the two main regions so identified, there is often evidence of the presence of a third located right at the surface, which is indicated to be amorphous, having cooled too rapidly to form crystals. In some studies, more than three distinct regions have been found.2'*22More references to the skin-core morphology in unfilled polymers are listed by Chen and White.23 Reinforcement can influence the morphology in three different ways. Firstly, the fibres or particulates modify the flow properties of the polymer and as a consequence may modify the morphology of the polymer itself. Secondly, in the case of crystallizing polymers, the additive may act as a crystal nucleant, influencing the crystal size and (local) orientation. Finally the distribution of the additive itself within the moulding (if segregation occurs) and the variation of its orientation distribution (if the additive is not equiaxed) has a marked influence over the morphology and properties. Sometimes a skin can be identified that corresponds closely to the skin in the unfilled polymer. In the case of fibre reinforced injection mouldings a stratified layer morphology is found to occur. Various characteristic layer types have been observed including (i) regions in which the preferred orientation of the fibres is parallel to the melt flow direction; (ii) regions in which the preferred orientation is transverse to the melt flow direction; (iii) regions in which the fibres tend to lie in planes parallel to a surface of the moulding but do not display any preferred orientation within the plane; (iv) regions in which the fibre

Short fibre filled thermoplastics

27

orientation is random in three dimensions. This is reviewed below. Plate-like fillers such as mica and talc can show similar features but have not been studied as extensively.

2.3.2.2 Observation of the skin-core morphology The skin-core morphology in glassy polymers is generally not visible to the naked eye even when the moulding is sectioned. Fracture surfaces often show two or more distinct regions and the skin-core boundary may sometimes be visible, but its identification as such is not possible without corroborative evidence. This is because changes in fracture mechanism are influenced by the nature of the applied stress and by the geometry of the moulding as well as by changes in morphology. The most convenient method of investigating the variation in molecular orientation is to observe the moulding between crossed polars when the corresponding variation in birefringence causes the formation of dark and bright fringes in monochromatic light or coloured fringes in white light. If a section of the moulding is made, the fringes are close together in the skin whereas the interior of the moulding shows almost uniform birefringence or perhaps a gradual change. The method can be made quantitative by measuring the relative retardation through the thickness of the moulding, either using a compensator or by identifying the colour produced when using white light, then machining away a thin layer and measuring the relative retardation of the remainder. This process is repeated until the remainder is too thin to handle. A plot of relative retardation versus depth removed is then generated and the birefringence at any chosen depth is given by the gradient of this plot.24-27The source of the birefringence in an injection moulding treated in this manner will normally be molecular orientation, though residual moulding stresses will also contribute. The problem of separating the two contributions is discussed e l ~ e w h e r e . ~ ~ - ~ ' Although semi-crystalline polymers are birefringent the observation of fringe patterns when viewed between crossed polars is only possible with thin sections because of the light scattering that occurs as the result of the presence of two phases (amorphous and crystalline) of very different refractive indices. Thin sections are normally prepared using a microtome and the transition between the spherulitic core and the oriented skin is usually very clear when the section is viewed between crossed polars in a low magnification microscope. The skin-core boundary is often visible on fracture surfaces, both in the light microscope and in the scanning electron microscope. The most important morphological variation is crystal orientation and is best investigated by X-ray techniques. Thin sections can again be used to isolate a chosen region. For the most comprehensive analysis, pole should be constructed for several prominent crystal direct ions . With filled polymers, light microscopy of a polished section is again a valuable technique, but the preparation of satisfactory sections is difficult because of the vast difference in hardness between the polymer and the filler. Thin sections for transmission microscopy are particularly difficult to make. The layer structure is normally well defined once a satisfactory section has been prepared.

28

Short fibrepolymer composites

2.2 Fibre cross-sections at different depths of a glass reinforced nylon 6,6 moulding: (a) near to the surface, where fibres are predominantly parallel to the flow direction and lie mainly normal to the section, presenting nearly circular profiles; (b) in the interior in a region in which many fibres lie at oblique angles to the section. (Micrographs courtesy of B O’Donnell.)

2.3.2.3 Fibre orientation distribution

The fibre orientation distribution has a strong influence over the properties of the material. The reinforcement provided by each individual fibre depends on the orientation with respect to the loading axis, while electrical and thermal properties are likewise affected. A short review of theoretical predictions of fibre orientation distribution is-given in Chapter 1, section 1.3. These are required for computer modelling in which the objective is to design the mould geometry to give a component with the target properties. For the purposes of testing the validity of theoretical predictions of fibre orientation distribution and of the

Short fibre filled thermoplastics

29

corresponding composite properties, measurements of the fibre orientation distribution are required. Ideally, it is necessary to determine the three dimensional orientation distribution of the fibres as a function of depth within the moulding. The orientation of a single fibre can be determined from the ratio of the minor to major axis of its elliptical profile in the section, Fig. 2.2. This equals the cosine of the angle between the fibre axis and the normal to the plane of the section. For sections cut perpendicular to the flow direction in the moulding this is the direction of paramount importance; the directions relative to axes chosen within the section plane can be measured directly using the major axis of the fibre section as reference. The fibre orientation distribution at a particular depth can be obtained by sampling a region at the selected depth (say a band from 0.5-0.6mm from the surface of the moulding) and determining the orientation of each of the fibres within it.32 This can be done using image analysis equipment, as discussed in Chapter 1. The results can be presented as histograms of the frequency of occurrence of each band of orientation, say 1&20", 2&30° from the flow direction.33 This process must be conducted for all depths within the moulding. Examples of histograms obtained from a glass fibre reinforced nylon 6,6 moulding are shown in Fig. 2.3 and confirm that axial orientation is favoured near the surface whereas the fibres are predominantly transversely oriented in the centre.33An alternative method that gives the projection of the fibre orientation is to prepare a thin section and form a shadow image of the fibres using X-rays ('mi~roradiography')~"~~ (Fig. 2,4). Another non-quantitative method for assessing orientation distribution is by scanning electron microscope examination of fracture (Fig. 2.5). The large depth of field of the scanning electron microscope tends to exaggerate the inclination of the fibres to the viewing direction and it is necessary to analyse images taken with the sample tilted at different angles to obtain even a semi-quantitative assessment of orientation distribution. In Fig. 2.5(b) there are several channels in the fracture surface corresponding to fibres that must have been removed, attached to the other fracture surface. These channels show unambiguously that the fibres were parallel to the fracture surface. Although the fibre orientation distribution has the greatest effect and has received most attention, the fibre volume fraction is an important parameter and investigations have been made to determine whether any segregation of fibres Kubat and Szalanczi used a spiral test mould and showed that there was a slight tendency for the fibre concentration to be greatest near the end of the flow path.39 Hegler and Mennig used a more conventional tensile test bar mould and found negligible variation in fibre concentration along the flow path or through the thickness of the moulding.40Both studies revealed that glass spheres segregated much more than fibres both along the flow path and through the thickness. More recently, Akay and Barkley found that the through-thickness segregation of glass fibres in tensile bar mouldings made from polypropylene and polyamide composites was small when short fibres were used but that consider-

Short fibre-polyrner composites

30 40

30 v)

Q) L

n .-

rc

20

1

10

0 5

15

25

35

45

55

65

75

a5

d 40

30

20 P rc bQ

10

I

0

5

15

25

35

45

#

55

I

I 65

75

a5

31

Short fibre filled thermoplastics

40

30 v)

z

e

20

w-

bQ

10

(4

5

15

25

35

45

55

65

75

# 2.3 Histograms giving fibre orientation distributions at different depths within glass fibre reinforced nylon 6,6 mouldings for intervals of 4 of lo", centred on (p = 5", 15 ', etc: (a) surface layer 0.18mm deep; (b) layer between depths 0.62-0.88 mm; (c) 1.55-1.81 mm (spanning the m i d - ~ l a n e ) . ~ ~

able segregation occurred with long fibres, the core fibre content exceeding that near the surface by about 2.3.2.4 Crystal nucleation effects Crystallization is often nucleated by fibres present in crystallizing thermoplastics. Nucleating agents are often added to moulding resins to produce smaller, more uniform, crystals, which usually improves the strength of moulded products. Nucleation on reinforcing fibres may give similar benefits but leads to a more complex morphology. The crystal orientation is strongly influenced by the surface on which nucleation takes place and the characteristic morphology in a crystallizing polymer matrix surrounding fibre nuclei has been called transcrystall in it^^^-^^ (Fig. 2.6). This does not necessarily result in columnar Nucleation may not occur simply as the result of surface interaction but may require, in addition, molecular orientation produced by flow during a moulding operation. In this case, transcrystallinity may be obtained in fibre-polymer combinations that do not show nucleation when s t a t i ~ . ~ 'It- ~is~generally considered that this interaction between the fibres and the matrix will improve the adhesion. The transcrystalline material has different properties from

32

Short fibre-polymer composites

2.4 Microradiographs of sections cut from glass fibre reinforced polystyrene: (a) cross-section through a bar 3 mm thick; (b) part of a section in which the initial fibre images were magnified using diverging X-ray source produced by an electron beam focused on a target in a scanning electron micro~cope.~' (Photographs courtesy of A K Srivastava.)

spherulitic material crystallized from the same polymer.50 In other studies, Burton et a1 present evidence for chain-extended structures growing from carbon fibre ends in reinforced p~l ypr opylene.Tan ~ ' et al observed nucleation at the crossover points of carbon fibres in their compression moulded samples.52 Pate1 and Bogue observed enhanced orientation of molecules in extruded non-crystalline polystyrene in the presence of glass fibres.53 Injection moulding also causes strong preferred orientation with platelet fillers such as talc or When crystals nucleate on the surfaces of the platelets then the crystal orientation distribution is strongly dependent on the particle orientation d i ~ t r i but ion.~ 1 * 5 7 - 5 9 In the case of platelets, the normal to the surface on which crystals nucleate has a unique preferred direction whereas with fibres the normal to the surface on which crystals nucleate is cylindrically distributed. As a result, the influence of oriented platelets on the orientation distribution of the polymer crystals is more marked than that of fibres.

2.3.2.5 Effect

of moulding conditions

Bright et a1 observed that the orientation of fibres in the core was dependent on the injection speed.36 High injection speeds produced strong fibre alignment transverse to the flow direction whereas the preferred orientation was parallel to flow at low injection speeds.36 This may be related to the observation that

Short fibre filled thermoplastics

33

Scanning electron micrographs of fracture surfaces taken from glass fibre reinforced polypropylene mouldings broken in tensile tests: (a) single end-gated bar; (b) double end-gated bar, showing a greater tendency for fibres to lie parallel to the fracture surface (and therefore perpendicular to the bar axis).

2.5

addition of fibres caused a significant increase in viscosity at low shear rates, whereas at high shear rates fibre filled resins had viscosities similar to unfilled resin of the same kind.36 In a later study Bright and Darlington found that the overall fibre orientation distribution in simple mouldings made from glass fibre reinforced polypropylene and nylon was not very sensitive to moulding conditions, including different mould temperatures.60 In a majority of products it is generally beneficial to maximize the degree of fibre orientation in the flow direction. This is especially difficult to achieve in thick mouldings but Bevis and co-workers have developed a method of live-feed

34

Short fibrepolymer composites

2.6 Light optical micrograph showing the interfacial microstructure for a poly(ethy1ene terephthalate) fibre in polypropylene having a melt flow index of 22. (Photograph courtesy of M J F ~ l k e s . ~ ~ )

injection in which two or more injection cylinders work in push-pull fashion, sending the melt repeatedly through the mould and depositing a layer of material at each pass (“SCORIM’process).6’-62This procedure is capable of producing very thick mouldings free from voids and, in the case of reinforced resins, with a high degree of fibre orientation. High orientation is present in some regions in mouldings made from ‘long fibre’ resins, as expected, but there are also zones in which swirl patterns prevail. Interestingly, fibre bundles formed during the manufacture of the moulding granules often remain intact when processing, and fibre length degradation is not as serious as might have been expected.63 If a moulding made from a ‘long fibre’ resin is ashed a fibre skeleton remains which retains the shape of the original article as a result of the high proportion of long fibres that survive the moulding process.64

2.3.2.6 Weld lines Knit lines or weld lines occur in injection mouldings when two melt fronts merge.64-68This happens when material enters the mould from more than one gate or on the downstream side of an obstruction such as an insert around which the melt flows. In an unfilled polymer a successful weld requires intimate mixing of the molecules across the knit line and this may be achieved if the temperature remains sufficiently high for an extended period. Any property advantage derived from molecular orientation may be lost during this process, however, and it is

Short fibre filled thermoplastics

35

important to design the placement of gates and inserts such that knit lines are avoided in critical load-bearing locations. In the case of reinforced resins the fibre orientation distribution at the knit lineis completely different from that in regions of uninterrupted flow. In the extreme example of the weld formed when two flow fronts travelling in opposite directions meet, the fibres tend to lie transverse to the flow direction at the weld (see Fig. 2.5(b))and the property enhancement due to fibre alignment is completely lost. Even if good adhesion across the weld were to be achieved the weakness caused by the loss of fibre alignment would normally be unacceptable. Studies of the characteristics of such welds in tensile test bars made in moulds gated at both ends are quite common. Savadori et UP*found that weld quality depended on the size and shape of the filler and that welds made when the filler was in the form of glass spheres were far superior to those made with glass fibre reinforcement.

2.4

Properties

2.4. I

Introduction

In this section the overall properties of injection mouldings made from fibre filled thermoplastics are considered. From the point of view of both the moulder and the end user the most important properties are the dimensions and the mechanical properties and these are the ones dealt with here. 2.4.2

Shrinkage

Several types of shrinkage are encountered in injection moulding and each of them will be considered here in turn. The most commonly acknowledged form of shrinkage is that which occurs when the moulding cools down to ambient temperature at the end of the moulding operation. This is often called mould shrinkage and a measure of the reduction in linear dimension that occurs is normally included on the data sheet supplied by the moulding resin manufacturer. The shrinkage is expressed as the fractional difference in dimension between the moulding and the mould cavity. This is the conventional way to describe shrinkage because it is simple to measure and is the most convenient form for a designer to use, but shrinkage is essentially a volumetric property (though it need not be isotropic). A further change in the dimensions of the moulding may take place over an extended period of time (months or even years) even if the moulding is kept at modest temperatures throughout. This is an example of physical ageing,69and is, of course, related to changes in density (see also section 2.5.2). If the temperature is raised then the moulding is likely to warp as a consequence of differential relaxation of residual stresses and/or differential re-coiling of molecules that were frozen into extended conformations when the moulding cooled. Differential re-coil may occur because the extent of molecular orientation varies from one location to another, particularly through the thickness of the moulding. Thus shrinkage may be used to give information about

36

Short fibre-polymer composites

the orientation state at different positions within the moulding, preferably by studying small samples removed from the chosen locations and heated in a controlled manner as described by Chen and White.23

2.4.2.I Mould shrinkage Material is delivered to the mould at high pressure, usually in the range 3@250 MPa, and at the completion of the filling phase the mould is overpacked. Cooling commences and the polymer, most of which is still molten, proceeds to contract. As long as the gate into the mould cavity remains open and pressure is applied by the injection system, more material is forced into the mould to counteract this shrinkage and to minimize voiding and sinking (‘topping-up’). If the dimensions of the gate are generous it may be possible to overpack the mould to such an extent that it remains overpacked even after cooling the material sufficiently to eject it from the mould. If this is the case the moulding will stick unless the cavity is shallow and tapered. Normally the gate is of modest dimensions and it freezes off quite early on in the moulding cycle, so preventing any further topping-up. If the mould cavity is much deeper than the gate then the material in the interior of the moulding may be at a much higher temperature when this happens and will subsequently undergo substantial thermal contraction. Freeze-off occurs at a higher temperature if the pressure is increased with both glassy and semi-crystalline polymers, as is evident on inspecting the volumetric data presented in the book by Mills70 and the original sources referenced therein. Some guidance on how to predict the amount of shrinkage in a moulding from the thermal properties of the polymer is given in the book by McCrum et a1.’l Before dealing specifically with fibre filled polymers it is necessary to consider the shrinkage characteristics of the unfilled polymers. 2.4.2. I. I Glassy polymers

The volumetric shrinkage behaviour of a glassy polymer is shown schematically in Fig. 2.7. The specific volume ( = l/density) is plotted as a function of temperature over a range that includes the glass transition temperature, T,. The slope of the volume-temperature plot above T, is much greater than that below it; the slope below T, is somewhat dependent on the cooling rate. Also shown in Fig. 2.7 are the temperatures T , at the time at which the gate freezes and T,, the final (ambient) temperature. Thus the shrinkage of the material at a given location is that suffered as the temperature falls from T , to T,. Account should be taken of the difference in pressure of the material at T , (equal to the holding pressure exerted by the injection system) and at T , (atmospheric pressure). At the time that the gate freezes there will be quite large temperature variations present within the moulding so that the value of T , will differ from one location to another. T , also depends on (i) the processing temperature, because the cooling characteristics at the centre of the gate and at any other chosen location are not linearly related; (ii) the holding pressure; and (iii) the dimensions of the gate. Mould shrinkage values quoted in data sheets can therefore be taken only as guidance. Shrinkage must be

Short fibre filled thermoplastics

37

/

glassy polymer

/

T2

T,

Tl

temperature 2.7 Schematic of the relationship between specific volume and temperature

(I)for a glassy polymer. The slope of the solid line below T, depends on the cooling rate. The broken line represents the equilibrium volume for temperatures below T,. T , is the temperature of the material at a given location within a moulding when the gate freezes. The value of T , depends on the location, T, is the ambient temperature; all parts of the moulding reach this temperature eventually.

allowed for when designing a mould and, for tight dimensional reproducibility, the processing parameters must be kept constant. 2.4.2. I .2 Semi-crystalline polymers

Shrinkage in mouldings made from semi-crystalline polymers is generally much greater than that with glassy polymers. This follows from the large decrease in volume that usually takes place on crystallization. The regular packing of molecules in the crystal phase usually results in a much higher density than is possible in the disordered phase, though there are exceptions, such as poly(-4 methyl pentene-1) in which the loose helical conformation of the molecules in the crystal phase causes it to have a density almost equal to that of the amorphous phase. More common is the behaviour shown by polyethylene in which the crystal density is 998 kgm- and the amorphous density is around 854 kgm-3, giving a volumetric reduction of nearly 17% in any region that crystallizes completely. Other crystal and amorphous density values have been tabulated by Gei1.72

38

Short fibrepolymer composites

I I I

T,

I I

I I

I I

I

I

I

T,

T1

T2

temperature 2.8 Schematic of the relationship between specific volume and temperature (gfor a semi-crystallinepolymer. The temperatures T , and T2 and the lines below T, have the same significance as in Fig. 2.7. T, is the crystallization

temperature.

As with glassy polymers, the critical event is the freezing ofthe gate. Some of the material near the mould cavity wall will be solidified at this time, but in the interior much of it will still be molten. The subsequent volumetric behaviour of such material can be followed in Fig. 2.8. Suppose the element of material under examination is at temperature TI at the time that the gate freezes. Cooling produces a fairly steep volumetric decrease until the crystallization temperature T, is reached. Then commences a very rapid fall in volume while crystallization proceeds. Below T, both the crystalline and the amorphous phases display volumetric contraction as the temperature continues to fall, but there is an additional contribution. Because the cooling rates prevalent in injection moulding are fast, some crystallization may continue after the temperature falls below T,, and the gradient of the volume-temperaturecharacteristic will differ from one position within the cavity to another according to the local cooling rate as well as the degree of crystallinity. After reaching T, the rate of volumetric contraction reduces, reflecting the behaviour of the amorphous phase and also because further crystallization will effectively cease. For many semi-crystalline polymers T, is below ambient temperature, T 2 , as in the example shown in Fig. 2.8. Whether T , falls above or below T,, the overall volumetric reduction in a

Short fibre filled thermoplastics

39

moulding made from a semi-crystalline pdymer is likely to be quite large, and so, therefore, is the linear mould shrinkage. 2.4.2. I .3

Thermotropic liquid crystal polymers

Only a small number of thermotropic liquid crystal polymer injection moulding resins have been made available commercially at the time of writing and only limited data are published on this family of materials. Nevertheless it seems that the shrinkage behaviour is likely to be fairly similar for all such polymers. Unlike the amorphous and semi-crystalline polymers discussed above it is no longer appropriate to discuss shrinkage in volumetric terms. Liquid crystal polymer mouldings have a highly developed anisotropy and the thermal expansion coefficient may vary strongly with direction. The presence of fibres may even enhance the polymer orientation still further.73 The thermal expansion coefficient of rigid rod units is very small parallel to the rigid rod axis (and may even be negative in this direction) so that the linear shrinkage of injection moulded liquid crystal polymer is very small when such units are oriented by flow. The volume shrinkage is much less than that of other thermoplastic materials. This is one of the key commercial properties of this class of polymer. 2.4.2. I .4

Filled thermoplastics

The shrinkage behaviour of filled polymers must be considered with reference to unfilled polymers, as discussed above. One of the motives for using filled polymers for engineering applications is that they have better dimensional properties than the corresponding unfilled resins. Filled thermoplastics display lower mould shrinkage because the filler particles in common use (glass, carbon, chalk, talc, mica, etc) have much lower thermal expansion coefficients than the host polymer, so diluting the effects of the shrinkage that takes place in the polymeric phase. Another possible effect is that the filler may provide nuclei for crystallization in which case crystallization is likely to be more nearly complete when the gate freezes, reducing the contribution to shrinkage that is caused by crystallization. Shrinkage is anisotropic if there is preferred fibre orientation. In injection mouldings possessing strata with different orientation distributions this can lead to warping if the layers are not symmetrically balanced across the centreline. Particulate reinforcement does not have this problem. 2.4.2.2

Shrinkage on ageing

The changes in linear dimensions that occur at room temperature are small even at long ageing times and are difficult to measure with sufficient accuracy to follow the process closely. It is relatively easy to measure the changes in density under these conditions, however, and these are discussed in section 2.5.2. 2.4.2.3

Shrinkage at elevated temperature

When the temperature of a moulding is raised to close to the softening temperature some shrinkage occurs. There is usually greater shrinkage in the flow direction than transverse to it, indicating that a significant part of the shrinkage is caused by re-coil of molecules frozen into extended conformations. Thus a shape

40

Short fibre-polymer composites

2.5

h

z x

----

GFPP

2.0

Y

1.5 U Q

-0

1.o

0.5 0 0

10

20

40

50

crosshead displacement

(mm)

30

60

2.9 Load-deformation curves for polypropylene and glass fibre reinforced polypropylene. The polypropylene sample continued to draw to several hundred millimetres. (Courtesy of Li Tong.)

change occurs. Additionally the moulding shows distortion of the kind that would be expected if shrinkage in regions near the surface were greater than in the interior, consistent with the variation in molecular orientation present through the thickness. This is discussed further in section 2.5.3 where the results of measurements made on samples extracted from different locations within mouldings are considered. 2.4.3

Mechanical properties

Fibre reinforcement improves both the stiffness and the strength of thermoplastics. Comparison of the load-deformation behaviour of unfilled polypropylene (PP) and glass fibre reinforced polypropylene in a standard tensile test at room temperature reveals that the reinforced material is both stiffer and stronger and that reinforcement inhibits drawing (Fig. 2.9). In some applications, the most useful benefit provided by fibre reinforcement is a very considerable increase in the creep resistance. If there is a marked preferred orientation of fibre alignment then the creep resistance is anisotropic and studies of this phenomenon have tended to distract from the overall improvement in creep properties that occur on adding fibres to thermoplastics.' Numerical

Short fibre filled thermoplastics

41

comparisons are difficult because of the non-linear, time-dependent nature of polymer creep and also because reinforced grades are stiffer, leading to further problems in choosing a basis for comparison - should it be for the same applied stress, or the stress to produce the same instantaneous strain, or a given fraction of the ultimate tensile stress, or some other criterion? Nevertheless the resistance to creep can be improved by orders of magnitude by introducing reinforcement into thermoplastics. A comparison of the properties of short fibre and long fibre reinforced polypropylene and polyamide injection moulded bars has been presented by Akay and B a r k l e ~ . ~They ’ varied the skin thickness by altering the moulding conditions and showed that both the stiffness and the strength depended on the skin
Sometimes the addition of fibres improves the toughness as well as the stiffness and strength, though in the case of thermoplastics that are tough when unfilled, the effect of adding filler is often to reduce toughness. The properties are generally dependent on the fibre-matrix adhesion, and Young’s modulus and Poisson’s ratio are reported to fall sharply if de-bonding occurs in nylon &glass bead composite^.^^ Injection moulded bars made from unfilled polypropylene do not break when bent by hand but similar mouldings made from glass fibre reinforced polypropylene break quite easily. This is because the presence of the fibres restricts the amount of rearrangement that can be achieved by the molecules of the polymer in response to an applied deformation. In a tensile test unfilled polypropylene bars neck and cold draw to several times their original length whereas glass fibre reinforced polypropylene bars break without necking (Fig. 2.9). It is difficult to predict whether an unfilled or a fibre reinforced grade will be superior in an impact test (e.g. Izod or Charpy). This is because the increase in stiffness will produce higher energy absorption for a given displacement of the sample as it bends (before breaking), whereas fracture will occur at a smaller displacement than with an unfilled sample because of the greater brittleness. With some compositions these effects almost balance out. Friedrich has examined the factors controlling the toughness in some detail and has identified those cases in which an increase in the fibre volume fraction will improve the toughness and those in which it will cause a reduction in to~ghness.~’ The mechanism of toughening has been a subject of much study, in common with other fibre reinforced composites. When a growing crack encounters a fibre it tends to travel along or close to the fibre-matrix boundary rather than straight

42

Short fibre-polymer composites

through the fibre.3 The fibre breaks at its weakest point and pulls out of the matrix as the material on either side of the crack parts. Considerable energy is expended against the frictional resistance to this and accounts for a large part of the toughening. The toughness of the composite is partly dependent on the fibre-matrix adhesion. In some cases low fibre-matrix adhesion gives highest toughness because this favours fibre pull-out rather than fibre fracture near to the local crack surface. The frictional resistance is maintained by the residual stresses in the matrix around the fibre caused by differential shrinkage during cooling at the end of the moulding cycle. The contribution of fibre pull-out to toughening is rate-dependent. Bailey and Bader found that with nylon 6,6 reinforced with glass fibre the matrix acted in a brittle fashion in a Charpy test (fast fracture) and that some matrix adhered to the fibres, whereas in slow bending the matrix was ductile, and none adhered to the fibres.76Jang and Lieu also found that the crack propagation mode was highly r a t e - d e ~ e n d e n t .Fibre ~ ~ avoidance dominated at low rates whereas the crack path was flatter at higher rates with less de-bonding and pull-out. The fracture toughness decreased at higher rates despite increased fibre fracture.77On the other hand, in fatigue tests on short and long fibre reinforced nylon 6,6, Karger-Kocsis and Friedrich observed increased ductility with increased stress intensity factor range and proposed that this was due to crack tip heating.78 Crack initiation is also related to the presence of the fibres. Sato et ul observed deformation of glass fibre reinforced nylon 6,6 in three-point bending in the scanning electron microscope and noted cracks formed at the fibre ends first, then along the interface, then in the matrix.79 Weiss" observed that failure in short graphite fibre-polypropylene mouldings was by a combination of matrix fracture and de-bonding while Yuan et UP' noted that these were competitive processes in their short fibre composites with poly(viny1chloride) matrix. These observations are consistent with the predictions of the mechanical analysis presented in Chapter 1 (section 1.4).

2.4.5 Thermal properties As with mechanical properties, the linear thermal expansion coefficient of a fibre reinforced composite is very anisotropic if the fibres have a preferred orientation. The expansion coefficient parallel to the fibres is much smaller than that in the transverse directions. The thermal conductivity is also changed by the presence of fillers, which generally have higher conductivities than the host polymers, giving increased rule of mixtures values in the composite. Even glass fibres have significantly higher thermal conductivity than the common host polymers. The thermal conductivity of carbon fibres is much higher than that of the polymers, and it is anisotropic.

2.4.6 Resistance to weathering and other environmental attack Most thermoplastics degrade rapidly when in outdoor service because of photo-oxidation promoted by ultraviolet (UV) irradiation.82 The reaction is

Short fibre filled thermoplastics

43

2.10 Surface of glass fibre reinforced polypropylene after outdoor exposure in Jeddah, Saudi Arabia, for one year. (Exposure conducted by M M Qayyum.)

fairly rapid in tropical regions (and under accelerated artificial weathering conditions) and oxygen is used up in the reaction before it can diffuse to the interior so that degradation is concentrated near the ~urface,’~*’~ even in polymers in which high UV levels are present in the interior.” In polymers that are fairly transparent to UV this effect is reduced if oxygen levels recover during dark periods, allowing some limited photo-oxidation in the interior when UV illumination recommences. In reinforced polymers the presence of the filler causes the UV intensity to fall to negligible levels at depths beyond typically 0.8mm of the ~urface.~’ Degradation deeper in the interior is consequently very limited and is not much affected by the introduction of a dark period. The surface appearance can be badly damaged when the ultraviolet intensity is high and fibres lying near to the surface can become exposed (Fig. 2.10). Significant reduction in strength occurs in samples around 3 mm thick as a result of the surface degradati~n.~‘.~’ The combination of glass fibre filler and carbon black was found to give good weather r e ~ i s t a n c e . ~ ~ There are several ways in which the effect of an aggressive environment on a fibre reinforced thermoplastic may differ from that on the unfilled polymer. A penetrant such as a cracking agent may diffuse along fibre-matrix boundaries more readily than through the matrix itself, resulting in degradation in the interior much sooner than in unfilled polymer. The penetrant may assist de-bonding at the fibre-matrix interface. The penetrant may promote fibre degradation. The latter was observed by Lhymn and SchultzE8 when they exposed poly(ethy1ene terephthalate) to 10% hydrochloric acid. This caused the fibres to fragment and, as a consequence of the loss of the pull-out contribution, the toughness decreased.

44

Short fibrepolymer composites

2.5 Effect of fabrication on properties 2.5./

lntroduction

Th overall prc erties of injection moulded fibre filled thermoplastics were dealt with in section 2.4 but it is evident that the properties are to some extent dependent on the moulding conditions. This is because the moulding conditions influence the layer structure, and the different layers have different properties. It is not yet known to what extent the layer structure can be controlled, but this section reviews the property variations through the depth of mouldings. This gives an indication of the limiting properties that might be achieved if full control over layer morphology were possible. It also provides information that relates to problems that are a direct consequence of the layer structure. For example, at the interface between layers, the mismatch of properties such as Young's modulus and thermal expansion lead to the development of delamination forces if large deformations or temperature changes are applied.

2.5.2 Density The density of injection mouldings varies with time and from one location to another. During the moulding process the material cools rapidly, even in the core in thin-walled mouldings, and the solidified polymer is far from equilibrium. As a result the material may slowly densify over a period of time. This is a property shown by glassy polymers as well as semi-crystalline polymers, in which additional crystallization can take place long after the moulding has reached thermal equilibrium. It is noted also that the material at every location within a moulding has a different thermal history, leading to differences in density from one position to another in both semi-crystalline and amorphous polymers. Additionally the frozen-in molecular orientation is significant near the surface of the moulding but much less (often not measurable) near the centre; differences in density at different locations may arise if the molecules are packed closer when oriented. An example of the variations in density in injection moulded (unfilled) polystyrene with location and time is given by Iacopi and White.89The material extracted from the skin is considerably denser than that extracted from the core; there is probably a steep density gradient within the skin, though this is not easy to measure because of sampling difficulties. The density appears to be uniform, or nearly so, over the majority of the section (core). The material within both skin and core becomes denser on ageing at room temperature at a fairly similar rate, approximately proportional to log (time).89 With some semi-crystalline polymers it has been observed that in the as-moulded state the density is least in the skin both when ~ n f i l l e dand ~ ~ when .~~ filled.23*91-94 Ho uska and Brummelgofound that the minimum density occurred about 0.25 mm from the surface of unfilled polypropylene mouldings, coincident with the location of the minimum crystallinity, as indicated by infra red measurements. It seems that, at the very fast cooling rate near the surface, crystal

Short fibre filled thermoplastics

45

growth is restricted giving a lower than average crystallinity and causing correspondingly a low density. The non-crystallized polymer retains preferential alignment of the chain axis parallel to the flow direction in the solidified skin, as confirmed by the significant shrinkage obtained on heating at elevated temperat u r e ~ Although .~~ the crystallinity in the skin is lower than in the core, leaving greater scope for further crystallization during ageing, Chen and White found that at room temperature the denser core and intermediate layer in unfilled mouldings became denser more rapidly than the skin, increasing the margin of difference in density still f ~ r t h e r . ’A~ possible explanation has been offered in terms of the different states of molecular orientation of the molecules in the non-crystallized phase at different locations through the depth of the mouldi n g ~The . ~ presence ~ of glass fibres or chalk filler reduced the rate of densification and Chen and White speculate that this is a consequence of crystal nucleation by the filler, causing primary crystallization to occur at a higher temperature when the molecules have greater mobility and are less inhibited by entanglements; the material that is rejected from the crystal phase under these conditions is less likely to be available for subsequent secondary crystallization on ageing than if crystallization occurs at a lower temperat~re.’~ 2.5.3

Shrinkage at elevated temperature

Samples of unfilled and filled polypropylene were extracted from different depths within moulded bars and heated to temperatures in the range 15&170°C in a study by Chen and White.23The critical temperature at which the onset of large scale molecular recoil occurs is around 160°C and skin samples from newly moulded unfilled bars and bars aged for two weeks all showed considerable shrinkage (> 5%) when heated to 170 “C,but less than 1YOwhen heated to 150 “C. The shrinkage observed in the core and in the intermediate layer was much less than in the skin. The samples extracted from polypropylene filled with glass fibre showed behaviour basically similar to that observed in unfilled polypropylene, indicating that the presence of the fibres did not inhibit molecular orientation nor subsequent recoil at elevated temperature. 2.5.4 Young’s modulus

Young’s modulus varies with depth within fibre reinforced thermoplastic injection mouldings, corresponding to the layered variations in fibre orientation distribution. The variation of Young’s modulus through the depth of nylon 6,6 injection mouldings reinforced with glass fibre has been investigated by O’Donnell and White33 using layerwise fibre orientation distributions to calculate Young’s modulus. Figure 2.1 1 shows the effect of making calculations for a composite containing 33% by weight glass fibres based on two different fibre aspect ratios, giving upper and lower bounds (see section 1.4.2).The aspect ratios chosen (10 and 60) correspond to the ends of the range of values within which most fibres were observed to lie and the extent to which fibre length preservation

Short fibrepolymer composites

10

8

6

4

2

0 0

0.4

0.8

1.2

1.6

depth

2.0

2.4

2.8

3.2

(mrn)

2.11 Distribution of Young’s modulus through the depth of a glass fibre

reinforced nylon 6,6 moulding determined using the aggregate theory and based on fibre orientation distributions measured using image analysis.” The upper bound (B)was obtained using fibre aspect ratio /Jdr = 60 and the lower bound (0) was obtained using rf/dr = 10.

is beneficial is evident. The data in Fig. 2.1 1 were calculated using a depth-varying Young’s modulus for the matrix, taken to be the same as that measured in an unfilled bar of the same polymer moulded under the same conditions: this was the source of the sharp variation in the calculated modulus of the composite bar near to the surface. For comparison, the Young’s modulus distribution was also obtained by measuring the bending stiffness of bars then removing an equal thickness of material from both surfaces and repeating the procedure until the remainder was too thin to make satisfactory measurements, then using a difference method to work out the required The measurements made using this procedure are shown in Fig. 2.1 2 and compared with those in Fig. 2.11 by superposition in Fig. 2.13, which shows the through-depth modulus for approximately half of the bar thickness. The agreement between the direct measurements and those computed from the fibre orientation distribution is good, except near the surface. Young’s modulus was higher at the surface and considerably lower near the centre, where the prevalence of transverse fibre orientation limits the reinforcement. The magnitbde of Young’s modulus

Short fibre filled thermoplastics

47

2.12 Distribution of Young's modulus through half a bar (thickness 3 mrn) of the kind analysed in Fig. 2.11, determined from measurements of the

bending stiffness.33 obtained from the direct measurements was higher than that from the image analysis. The upturn at the surface was not apparent if the calculation based on the fibre orientation distribution used a uniform matrix Young's modulus instead of the variable and it was concluded that there is a depth-varying contribution to stiffness from the matrix phase even in a filled injection moulding. Unfilled mouldings show a strong depth d e p e n d e n ~ ebut ~ ~ this is normally overlooked in composite mechanics, where it is implicitly assumed that the fibre reinforcement will dominate. The average Young's modulus of the composite was approximately 9 GN m-' compared with approximately 3 GN m - 2 for the polymer resin. The high orientation produced by the SCORIM process (section 2.3.2.5) gives correspondingly high stiffness in the orientation dire~tion.~' Significant improvements in stiffness have also been obtained by orienting the fibres using die drawing and melt extrusion through a converging die.98 Although similar fibre orientation distributions were obtained by the two routes, the modulus of the die drawn material was much higher than that of the melt extruded composite. This was attributed to higher matrix orientation produced in die drawing, another example of controlling polymer matrix properties to achieve optimal properties.

48

Short fibrepolymer composites

20

I

'

"

'

"

"

'

15

v)

3 3

U

10

v)

oJ 5 c

J 0

> 0 0

1.o

0.5

depth

1.5

(mm)

2.13 Comparison of the Young's modulus distributions given in Fig. 2.1 1 and 2.12. The data points are those for the upper and lower bounds in Fig. 2.11 and the line is copied from Fig. 2.12.

2.5.5 Residual stresses Residual stresses form in injection mouldings as the consequence of the temperature gradients that exist during the cooling p h a ~ e . ' ~ Whe , ' ~ ~n the melt fills the mould, the material in contact with the cold mould cavity wall freezes instantly to form a solidified skin. The material in the interior cools much more slowly because of the low thermal conductivity of the polymer. As it cools it contracts, but pressure is held on after the mould is filled and material continues to flow into it from the injection system to keep it topped up. At some point the gate freezes and no more topping-up can occur. In all but the thinnest-walled mouldings this happens long before the material in the interior has cooled to the temperature of the mould and continued cooling results in the formation of residual tensile stresses in the interior, setting up opposing compressive stresses near the surface. The magnitude of the stresses and their detailed distribution depend on the moulding conditions and are most sensitive to the cooling rate. If the cooling rate is low then stresses are relieved by molecular relaxation processes almost as quickly as they form. Thus a low mould temperature favours high residual stresses. Mouldings are ejected long before the material in the interior reaches thermal equilibrium with the material near the surface and the cooling

Short fibre filled thermoplastics I

I

I

GFPP

iu I

E

5 -

49

.

- 0

--- - - - - -

0

0

-zI

-I

-

-

- ice

/

-10

I

-

- - _ _ - liquid N2

/' /

I

-

water

r

I' I

-1 5

0

I

I

I

0.5

1 .o

1.5

-

i 2.0

depth (mm) 2.14 Residual stress distributions in glass fibre reinforced polypropylene

bars cooled by ejecting theni prematurely from the mould and immediately placing them in ice water or liquid nitrogen. The bars were approximately 3 mm thick and the analyses are shown for half the bar thicknes~.''~

conditions of the moulding in the period immediately after ejection will also have a strong influence on the residual stresses. The magnitudes of the residual stresses are also sensitive to the presence of additive^.'^'-'^^ Residual tensile stresses as high as 20 MN m-' have been measured near the centre of injection moulded glass fibre reinforced polypropylene bars cooled in cold water immediately after ejection from the moulding machine, and then in liquid nitrogen. l o 3 Typical residual stress profiles in rapidly cooled mouldings made from glass fibre reinforced polypropylene are given in Fig. 2.14. Residual stresses can change markedly after moulding, particularly if the moulding is exposed to a high temperature or absorbs vapour or liquid from its surrounding^.'^ 1,103,105 The presence of fibres reduces the fractional changes in residual stress levels caused by ageing and heat treatment. l o 3 Residual stresses can have a strong influence over the distortion and the fracture of moulded polymers and steps may need to be taken to ensure that they do not contribute to failure. On the other hand, recent studies on unfilled thermoplastics have indicated that the compressive residual stresses near the surface of injection mouldings may retard photo-degradation and may therefore improve the weatherability. l o 4

50

Short fibre-polymer composites

The residual stresses discussed above are long range. In a composite there are in addition localized stresses around the fibre or filler particles, resulting from the differential shrinkage that takes place during post-moulding cooling between the fibre (which will generally have a quite small thermal expansion coefficient) and the matrix (which will generally have a quite large thermal expansion coefficient). This will cause the matrix to shrink onto the fibre, gripping it tightly even if there are no strong chemical bonds at the interface, and enhancing the frictional force on fibre pull-out (section 2.4.4).

2.6 Conclusions Short fibre reinforced thermoplastics have an attractive combination of properties and processability. Further improvements in property could be achieved if methods could be devised to ensure that the fibres were placed in the most advantageous manner. The use of particulate filler in conjunction with fibres may improve properties such as warpage resistance (for example, using a crystal nucleant such as talc in polypropylene) or weatherability (for example, using carbon black). References 1 Folkes M J, Short Fibre Reinforced Thermoplastics, Research Studies Press/Wiley, Chichesterrnew York (1982). 2 Franzen B, Klason C, Kubat J and Kitano T, Composites, 20 (1989) 65. 3 Mandell J F, Huang D D and McGarry F J, Shorr Fibre Reinforced Composite Materials, A S T M S T P , vol 772, Ed BA Sanders, ASTM, Philadelphia (1982) 3. 4 Shortall J B and Pennington D, Plust Rubber Proc Applics, 2 (1982) 33. 5 Bridge B, Folkes M J and Jahankhani H, J Mater Sci, 24 (1989) 1479. 6 Yang H H, Fibre Reinforcements for Composite Materials, Ed A R Bunsell, Elsevier, Amsterdam (1 988) 249. 7 Calundann G, Jaffe M, Jones R S and Yoon H Fibre Reinforcements for Composite Materials, Ed A R Bunsell, Elsevier, Amsterdam (1988) 21 1. 8 Allen S R, Filippov A G, Farris R J, Thomas E L, Wong C-P, Berry G C and Chenevey E C, Macromolecules, 14 (1981) 1135. 9 Acosta J L, Ojeda M C, Morales E and Linares A, J Appl Polym Sci, 31 (1986) 2351. 10 Morales E and White J R, J Mater Sci, 23 (1988) 4525. 11 Newman S and Meyer F J, Polyrn Compos, 1 (1980) 37. 12 Hawley G C , SPE 42nd A N T E C , New Orleans (1984). 13 Avella M, Martuscelli E, Sellitti C and Gargnani E, J Muter Sci, 22 (1987) 3185. 14 Chiang W-Y and Yang WD, J Appl Polym Sci, 35 (1988) 807. 15 Pape P G and Plueddemann E P , J Adhesion Sci Technol, 5 (1991) 831. 16 Garton A, Kim S W and Wiles D M, J Polym Sci Letts E d , 20 (1982) 273. 17 Kastner E, Nardin M, Papirer E and Riess G, J Muter Sci Lerts, 7 (1988) 955. 18 Denault J and Vu-Khanh T, Analytical and Testing M e t l d o l o g i e s f o r Design with Adunnced Materials, Eds G C Sih, J T Pindera and S V Hoa, Elsevier/North Holland, Amsterdam (1988) 141. 19 Bajaj P, Jha N K and Jha R K , Polym Eng Sci, 29 (1989) 557.

Short fibre filled thermoplastics

51

20 Arroyo Ramos M, Sanchez Berna M, and Vigo Mathem JP, Polym Eng Sci, 31 (1991) 245. 21 Bowman J, Harris N and Bevis M, J Muter Sci, 10 (1975) 63. 22 Bowman J and Bevis M, Plast Rubber Maters Applic, 1 (1976) 177. 23 Chen Z and White J R, Plast Rubber Compos Maters Applic, 18 (1992) 289. 24 White E FT, Murphy B M and Haward R N, J Polym Sci Polym Letts Ed, 7 (1969) 157. 25 Haworth B, Sandilands G J and White J R, Plast Rubber Int, 5 (1980) 109. 26 Cuckson I M, Haworth B, Sandilands G J and White J R, Intern J Polymeric Muter, 9 (1981) 21. 27 Campbell D and White J R, Polymer Characterization, Chapman and Hall, London (1989). 28 Read B E, Duncan J C and Meyer D E, Polym. Testing, 4 (1984) 143. 29 Saffel J R and Windle A H , J Appl Polym Sci, 25 (1980) 1117. 30 Isayev AI, Polym Eng Sci, 23 (1983) 271. 31 Chen Z , Finet M C, Liddell K, Thompson D P and White J R, J Appl Polym Sci, 46 (1992) 1429. 32 Fakirov S and Fakirova C, Polym Compos. 6 (1985) 41. 33 O’Donnell B and White J R, Plust Rubber Compos Proc Applies, 22 (1994) 69. 34 Darlington M W and McGinley P L, J Muter Sci, 10 (1975) 906. 35 Darlington M W, McGinley P L and Smith G R, J Muter Sci, 11 (1976) 877. 36 Bright PF, Crowson R J and Folkes M J , J Muter Sci, 13 (1978) 2497. 37 Hemsley D A and Hayles M, Developments in Electron Microscopy and Analysis 1977, (Instituteof PhysicsConference Series 36) Ed D L Misell, Institute of Physics, Bristol & London (1977) 53. 38 White JR and Thomas E L, Rubber Chem Technol, 57 (1984) 457. 39 Kuhat J and Szalanczi A, Polym Eng Sci, 14 (1974) 873. 40 Hegler R P and Mennig G, Polym Eng Sci, 25 (1985) 395. 41 Akay M and Barkley D, J Muter Sci, 26 (1991) 2731. 42 Burton R H and Folkes M J, Plast Rubber Proc Applies, 3 (1983) 129. 43 Burton R H and Folkes M J, Chapter 9 in Mechnnicul Properties of Reinforced Thermoplastics, Eds D W Clegg and A A Collyer, Elsevier Applied Science, London/ New York (1986) 269. 44 Folkes M J and Hardwick S T , J Muter Sci Letts, 3 (1984) 1071. 45 Campbell D and Qayyum M M, J Polym Sci Polym Phys Ed, 18 (1980) 83. 46 Campbell D and White J R , Angew Makromol Chem, 112 (1984) 61. 47 Devaux E and Chabert B, Polym Communs, 31 (1990) 39 1. 48 Thomason J L and Van Rooyen AA, J Muter Sci, 27 (1992) 889. 49 Varga J and Karger-Kocsis J, Compos Sci Technol, 48 (1993) 191. 50 Folkes M J and Hardwick S T, J Muter Sci Lefts, 6 (1987) 656. 51 Burton R H, Day T M and Folkes M J, Polym Communs, 25 (1984) 361. 52 Tan J K, Kitano T and Hatakeyama T, J Muter Sci, 25 (1990) 3380. 53 Patel P D and Bogue D C , Polym Eny Sci, 21 (1981) 449. 54 Crosby S M and White J R , Plast Rubber Proc Applics, 12 (1989) 171. 55 Xavier S F, Schultz J M and Friedrich K, J Muter Sci, 25 (1990) 241 1. 56 Xavier S F, Schultz J M and Friedrich K, J Muter Sci, 25 (1990) 2421. 57 Fujiyama M and Wakino T, J Appl Polym Sci, 42 (1991) 9. 58 Fujiyama M and Wakino T, J A p p l Polyrn Sci, 43 (1991) 97. 59 Fujiyama M, Intern Polym Processing V I l , 1 (1992) 84. 60 Bright P F and Darlington M W, Plasf Rubber Proc Applies, 1 (1981) 139.

52

Short fibre-polymer composites

Allan P S and Bevis M J, Plast Rubber Proc Applics, 7 (1987) 3. Gibson J R, Allan P S and Bevis M J, Plast Rubber Int, 16(5) (1991) 12. Gibson A G , Plast Rubber Proc Applics, 5 (1985) 95. Gore C F and Cuff G, 41st Annual Conferenceon Reinforced Plastics, Composites Inst SPI, Atlanta, (1986) Paper 8F. 65 Akay M and Barkley D, Plast Rubber Compos Proc Applics, 20 (1993) 137. 66 Haskell W E 111, Petrie S P and Lewis R W, Polym Compos, 4 (1983) 47. 67 Nadkarni V M and Ayodhya S R, Polym Eng Sci, 33 (1993) 358. 68 Savadori A, Pelliconi A and Rornanini D, Plast Rubber Proc Applics, 3 (1983) 215. 69 Struik L C E, Physical Aging in Amorphous Polymers And Other Materials, Elsevier, Amsterdam (1978). 70 Mills N J , Plastics, 2nd Ed, Edward Arnold, London (1993). 7 1 McCrurn N G, Buckley C P and Bucknall C B, Principles of Polymer Engineering, Oxford Science Publications, Oxford (1988). 72 Geil P H, Polymer Single Crystals, Wiley-Interscience, New York (1963). 73 Bharna S and Stupp P, Polym Eng Sci, 30 (1990) 228. 74 Van Hartingsveldt E A A and Van Aartsen J J, Polymer, 30 (1989) 1984. 75 Friedrich K, Compos Sci Technol, 22 (1985) 43. 76 Bailey R S and Bader M G, J Mater Sci Letts, 4 (1985) 843. 77 Jang B Z and Lieu Y K, J Appl Polym Sci, 30 (1985) 3925. 78 Karger-Kocsis J and Friedrich K, Composites, 19 (1988) 105. 79 Sato N, Kurauchi T, Sat0 S and Kamigaito 0, J Mater Sci Letts, 2 (1983) 188. 80 Weiss R A, Polyrn Compos, 2 (1981) 95. 81 Yuan J, Hiltner A, Baer E and Rahrig D, J Mater Sci, 20 (1985) 4377. 82 Davis A and Sims G, Weathering of Polymers, Applied Science, Barking (1983). 83 DeBruijn J C M, The Failure Behaviour of High Density of PoLyethylene with an Embrittled Surface Layer due to Weathering, Doctoral thesis, Delft University Press (1992). 84 Audouin L, DeBruijn J C M, Langlois V and Verdu J, J Mater Sci, 29 (1994) 569. 85 O'Donnell B and White J R, Polym Degrad Stab, 44 (1994) 21 1. 86 Qayyum M M and White J R, Plast Rubber Proc Applics, 12 (1989) 171. 87 Qayyum M M and White J R, Polym Compos, 11 (1990) 24. 88 Lhyrnn C and Schultz J M, J Mater Sci, 18 (1983) 2923. 89 Iacopi A V and White J R, J Appl Polym Sci, 33 (1987) 577. 90 Houska M and Brummel M, Polym Eng Sci, 27 (1987) 917. 91 Bakerdjian 2 and Kamal M R, Poiyrn Eng Sci, 17 (1977) 96. 92 Moy F H and Kamal M R, Polym Eng Sci, 20 (1980) 957. 93 Kamal M R and Moy F H, Chem Eng Commun, 12 (1981) 253. 94 Kamal M R and Moy F H, Polym Eng Revs, 2 (1983) 381. 95 ODonnell B, P h D Thesis, University of Newcastle upon Tyne (1990). 96 O'Donnell B and White J R , J Appl Polym Sci, 47 (1993) 189. 97 Hine P J , Duckett R A , Allen P S , Bevis M J and Ward I M , 3rd International Conference on Deformation of Fracture Composites, University of Surrey, March, IoM, London (1995) 536 . 98 Cook AH, Hine P J, Duckett R A and Ward I M, 3rd International Conference on Deformation of Fracture Composites, University of Surrey, March, IoM, London (1995) 526. 99 White J R, Polym Testing, 4 (1984) 165. 100 Isayev A I and Crouthamel D L, Potym Plast Technol Eng, 22 (1984) 177. 61 62 63 64

Short fibre filled thermoplastics

101 102 103 104 105

53

Thompson M and White J R, Polym Eng Sci, 24 (1984) 227. Morales E and White J R, J Mater Sci, 23 (1988) 3612. Hindle C S, White J R, Dawson D and Thomas K, Polym Eng Sci, 32 (1992) 157. Kwok CS, Li Tong and White JR, Polym Eng Sci, 36 (1996) 651. Paterson M W A and White J R, J Mater Sci, 27 (1992) 6229.

3 Thermosetting short fibre reinforced composites S B WlLKlNSON AND J R WHITE

3. I

Introduction

Thermosetting polymeric materials are produced by a number of short steps.'-2 At room temperature thermosetting resins are viscous liquids. The polymerization reaction, often called curing, is promoted under the application of heat. The product forms as a three dimensional network structure which cannot be melted. Common types of thermosetting polymers are polyester, phenolic and epoxy. Thermosets have a network structure, in contrast to thermoplastics which have a linear or branched chain structure. Thermosets cannot be easily shaped after polymerization and so are normally cured in moulds having the required finished shape. Special partly cured thermosets have been developed for use as matrices in composites. These are soft and easily mouldable and are cured by heating, usually while still in the mould. Reinforced thermosets generally have even better creep resistance than reinforced thermoplastics and usually permit higher service temperatures. Their disadvantages are that they are less convenient to store prior to moulding and that they are generally less easily fabricated than reinforced thermoplastics despite modern developments. When due attention is paid to the storage and fabrication conditions quite intricate shapes are possible, and reinforced thermosets are finding increasing applications (Fig. 3.1).

3.2 Curing characteristics Curing of thermosetting resins is very exothermic. This is most critical when thick sections are moulded. Addabbo et a13 studied the curing of a thermoset in a heated mould and showed the existence of a critical thickness below which the cycle time was not dependent on the part thickness. Williams et a14 found that the hottest plane does not always coincide with the centreline, and that the cure cycle time is not necessarily proportional to the part thickness, as often asserted. It is important to select the optimal wall temperature so that the mould can be filled without premature gelling, and so that the maximum temperature in the part stays below a ceiling at which degradation or undesired side reactions may occur.

Thermosetting short fibre reinforced composites

55

3.1 Car timing pulleys made from glass fibre-reinforced mineral-filled

phenolic. (Photograph provided by and published with the permission of Vyncolit NV, Ghent, Belgium.)

When the cycle time is determined by curing at the wall it is important not to overestimate the critical conversion needed to eject the part from the mould. The material must reach a critical conversion level at which it is dimensionally stable and can be removed from the mould without losing its shape or blemishing its s ~ r f a c eThis . ~ is often described as the end of cure, though there will normally be further reaction (‘curing’) after this time. Fillers absorb some of the heat and reduce the peak exotherm temperature. This improves the uniformity of the cure and should decrease residual stresses and lead to more accurate and consistent mouldings. It also reduces thermal degradation, allowing the use of higher mould temperatures, which may result in a reduced mould cycle time,6 maximizing efficiency. The cure time is sometimes taken to be the time to reach the maximum exotherm. To minimize the cure duration, the level of conversion throughout the mould needs to be accurately estimated or monitored. Three monitoring methods are in common use to determine the state of cure: 1 Stock temperature measurement, with completion of cure being indicated by a sharp rise in temperature; 2 Mould pressure measurement, with cure being indicated by a drop due to the rapid shrinkage of resin at cure; 3 Mould displacement measurement, where cure is indicated by a sharp decrease due to resin shrinkage at the cure point.

Another additive commonly used with thermosetting resins is a low profile additive, which is used to improve the surface quality of the moulding and prevent shrinkage. Low profile additives (LPAs) can form discrete particles or

56

Short fibre-polymer composites

co-continuous structures depending on the concentration. The addition of poly(viny1 acetate) as LPA to unsaturated polyester formulations is effective in ensuring that the moulded surface faithfully represents the surface of the mould cavity only when co-continuous structures are formed,. A critical amount of PVAc yields a co-continuous structure; adding more PVAc than this does not give any further improvement in surface quality. Although the effects of low profile additives and fillers on the cure kinetics have been the subject of several studies, the results are still controversial. Kubota7 found the effect of a low profile additive on cure kinetics was very small, whereas Han and co-workerss-12 reported lower reaction rates and lower final degrees of cure when using a low profile additive. Lem and Hans found that the effect of increasing the particulate filler content was to cause shear thinning and an increase in the rate of cure. The filler particles helped to control shrinkage during cure when the resin was subjected to steady shear deformation and the gel time was shorter for mixtures of resin and particles than for resin alone. An increase in the degree of crosslinking increases the glass transition t e m p e r a t ~ r e ’and ~ techniques have been developed to estimate the crosslink density by measuring the glass transition temperature. The dependence of properties on the composition and/or extent of cure is not easily generalized. Mostovoy and RepplingI4 found that ultimate strength was optimized by changing the ratio of curing agent to resin. Kim” concluded that tensile properties are nearly independent of the ratio of resin to curing agent. Yamini and Young16 observed a decrease in modulus with an increase in amount of curing agent beyond the theoretical stoichiometric point. Jonejai3 found that increases in Young’s modulus were attributable to additional curing, i.e. longer exposure to higher temperatures. Typical property changes promoted by post-curing polyester are increases in modulus of approximately 30% and in strength of around 60%.

3.3

Thermosetting resin types 3.3. I

Polyester resins

The general chemical structure for polyester resins is as shown below:

Polyester resins are commercially the most important group of resins for glass-resin composites. The resin is produced by a condensation reaction between glycol and an unsaturated dibasic acid. Both of these reagents contain carbon-carbon double bonds. Polymerization takes place through a radical chain reaction involving these double bonds and a vinyl monomer, usually styrene. In the case of composites the reaction takes place in the presence of the filler/fibrous phase. Further discussion of polyester resins is given elsewhere.’ 2 3

Thermosetting short fibre reinforced composites

57

Much of the versatility ofreinforced polyester systems lies in the wide variation in resin composition and fabrication methods possible, allowing the properties of the product to be tailored to the application required. The unsaturation in the polyester is usually supplied by the inclusion of maleic anhydride as one component. In addition a saturated acid or anhydride such as phthalic anhydride is often used. A higher proportion of unsaturated acid gives a more reactive resin with improved stiffness at high temperatures, whereas increasing the fraction of saturated components gives less exothermic cures and less stiff resins. Propylene glycol is the most widely used dihydroxy component of the polyester, and styrene is overwhelmingly the most popular crosslinking agent. A typical formulation for a general purpose polyester may contain 43 mass units of phthalic anhydride, 19 mass units of maleic anhydride and 38 mass units of propylene glycol. To each 100 mass units of this formulation is added 70 mass units of styrene for crosslinking. The amount of fibre reinforcement required for the resulting 170 mass units of resin will vary between 100 and 300 mass units. A form of reinforced polyester that has been developed for injection moulding is dough moulding compound (DMC)in which the fibres are bound together by unreacted monomer and which retains sufficient flow properties to process on reciprocating screw injection moulding machines. 3.3.1. I

Basic properties and applications of polyester resins

The most important properties of the unsaturated polyester systems include ease of handling, rapid curing with no volatiles evolved, light colour, dimensional stability and generally good physical and electrical proper tie^.^^ The major applications of fibre reinforced polyester resins fall into the following categories: boat hulls (including minesweepers and recreation craft); transportation, including bodies and other parts for passenger cars and truck cabs; consumer products, including such diverse items as luggage, chairs, fishing rods and trays; electrical appliances; construction applications, including sheet and panelling, pipes and ducts; and missile and radome uses. 3.3.2 Epoxy resins

The epoxy resin most widely used is made by condensing epichlorohydrin with diphenylol propane, more commonly known as ‘bisphenol An excess of epichlorohydrin is used to leave epoxy groups on each end of the low molecular weight polymer as shown overleaf. Epoxy resins are formed in two reaction stages. First a linear polymer is formed by the reaction between epichlorohydrin with a dialcohol (e.g. diphenylol propane) using aqueous sodium hydroxide as catalyst. It is in this condition that the resin is brought into contact with the fibres. The fibres have water adsorbed onto the surface, simply from exposure to air, and the ethylene oxide group contained in the resin reacts with the moisture, providing excellent adhesion. As distinct from polyesters, the linear polymers of this type cannot crosslink by themselves or with the assistance of a reaction initiator. Crosslinking occurs

58

Short fibre-polymer composites

instead by introducing chemicals which can react with the epoxide and hydroxy groups between adjacent chains and therefore the extent of crosslinking is governed by the quantities of curing agents added. Curing agents such as amines or acid anhydrides are added to 10-15% by weight and become part of the cured epoxide resin. The choice of curing agent and the proportion cured therefore contributes significantly to the properties of the cured epoxide resin.

3.3.2.I Basic properties and uses of epoxy resins The major use of the epoxy resins is as surface coating materials, which combine toughness, flexibility, adhesion and chemical resistance to a nearly unparalleled degree. Epoxy resins can be used in both laminating and moulding techniques to make fibre reinforced articles with better mechanical strength, chemical resistance and electrical insulating properties than those obtained with unsaturated polyesters. Only the higher price ofepoxies prevents their wider use within this field. Casting, potting, encapsulation and embedding are widely practised with the epoxy resins in the electrical and tooling industries. Other important uses include industrial flooring, adhesives and solders, foams, highway surfacing and patching materials. 3.3.3

Phenolic resins

Phenolic resins are among the oldest synthetic polymers available.25 The production of a phenolic resin is by a condensation reaction, which involves the evolution of water, and therefore special procedures for moulding the material are required. Phenols react with aldehydes to give condensation products if there are free positions on the benzene ring ortho and parm to the hydroxyl group. Formaldehyde is by far the most reactive aldehyde and is chosen almost exclusively in commercial production. The reaction is always catalysed by either acids or bases. The nature of the product is greatly dependent on the type of catalyst and the mole ratio of the reactants. The first step in the reaction is the formation of addition compounds known as methylol derivatives, the reaction taking place in the ortho and para positions (see top of next page). These products, which may be considered the monomers for subsequent polymerization, are formed most satisfactorily under neutral alkaline conditions.

Thermosetting short fibre reinforced composites

0

~!JCH~O-OCHK)H

OH +CH20-

59

OH

o OHc H 2 0 & CH2OH

CHPH

In the presence of acid catalysts and with the mole ratio of formaldehyde to phenol less than 1, the methylol derivatives condense with phenol to form first dihydroxy diphenyl methane: OH

OH

OH

OH

OH

OH

OH

OH

CH2OH A

t

H20

and on further condensation and methylene bridge formation, fusible and soluble linear low polymers called novolacs with the structure:

where ortho and puru links occur at random. Molecular weights may range as high as 1000, corresponding to ten phenyl residues. These materials d o not themselves react. 3.3.3.I Phenolic resin reinforced composites Unfilled phenolic resin has very few engineering uses. By including reinforcements and fillers, significantly better mechanical properties can be produced, giving increased reliability, better workability, reduced costs, better strength and improved impact resistance.26 3 ' The most popular fillers in phenolic resin composites are based upon finely divided wood flour from soft wood trees, which tends to improve impact resistance. A fibrous filler can be used in addition to improve mechanical properties. The fibrous phase is usually made from carbon, glass, aramid or silica, yielding a hybrid composite. The resin and the filler materials need to be mixed before moulding. The properties of the compound are determined by the size, concentration, aspect ratio, shape and length of the reinforcing fibres and on the interaction between filler and matrix. The fillers tend to increase the viscosity of the system and tend also to increase thermal conductivity which reduces the temperature rise during the exothermic curing reaction. These changes in viscosity are particularly important in the shaping/ moulding procedures. For example, when injection moulding phenolics, the higher viscosity of filled material causes higher injection pressures to be required,

60

Short fibre-polymer composites

increasing the running costs, but cycle times will be lower since the rate of heat transfer is improved, giving a better production rate. Another advantage of moulding reinforced phenolic materials is that composites can be designed with thinner sections owing to the superior stiffness and strength of the reinforced phenolic material. Disadvantages caused by introducing fillers are few and mostly insignificant, though some fillers can cause high wear in the processing machinery. With fibrous fillers orientation may occur and although this will generally produce beneficial properties, sometimes problems may occur, for example preferential orientation causes anisotropy and, in some cases, planes of weakness. High temperatures and excessive work during processing can cause separation of fibres from the matrix, though it has been found that by using large sectioned sprues and runners, the tendency for fibre degradation is reduced. Glass fibres are the most popular fibre reinforcement. The glass fibres give a hard impervious surface which is resistant to environmental corrosion by most media. The fibres also impart excellent high temperature resistance, high tensile strength and impact resistances up to 20 times that of an unreinforced resin sample. Although more conductive than the unfilled resin sample, the glass fibre reinforced phenolic materials still have low thermal and electrical conductivities compared with most other engineering materials and find application where good insulation properties are required. The optimum fibre content is 30-40% by weight.

3.3.3.2 Basic properties and applications of phenolic resins Phenolic resins when combined with suitable fillers and reinforcements have good chemical and thermal resistance, dielectric strength and dimensional stability.32 Products made with these resins are inherently low in flammability, are creep resistant and have low moisture absorption. Phenolic resins have excellent heat resistancecapabilities. They can be used as resins for products such as decorative laminates for counter tops. Phenolic resins are also used in engineering applications such as brake linings, abrasive wheels, sand paper and foundry moulds. Reinforced phenolic resins form a very creditable group of engineering materials with a wide range of applications covering a variety of industries. The impact grades are used in gearing, tool handles, washing machine parts, electrical switch gear, automotive parts, electrical controls and telephone parts. The automotive applications include flame-retardant linings with low smoke density when exposed to fire, brake pistons and high performance automotive transmission. The high fire retardance and low smoke density and toxicity features are also becoming recognized by many aircraft companies. The defence industry provides much initiative for the use of reinforced phenolic resins, particularly in research connected with missile technology where their relatively low density is valuable. Phenolic composites can be implemented as missile bodies, thrust chambers, rocket nozzles, motor casings and radome structure stiffeners. It is evident that modern phenolic materials can compete with other materials at the initial design

Thermosetting short fibre reinforced composites

61

stage and they need not be regarded simply as a cheap substitute for less critical components. Thus by the introduction of a fibrous reinforcement within a phenolic matrix the very attractive properties of phenolic polymer materials are upgraded to the extent where they are compatible in certain applications with the best of the exotic steels, refractory metals and other non-ferrous metal alloys.

3.4

Fabrication methods

Polymers can be moulded or shaped in a number of different ways. The choice of method depends on the type of material to be shaped (thermoplastic or thermoset, filled or unfilled); the cost of shaping the material; and the purpose of shaping the polymer, especially the application of the moulded part. The general definition of a moulding process is one in which polymer particles or melt are forced by the application of heat and pressure to flow into, fill and conform to the shape of a cavity (mould). The shaping processes reviewed here are all available for shaping short fibre reinforced polymer composites. 3.4. I

Compression moulding

In compression moulding the polymer is placed in a two-piece male/female mould.33 The mould is closed and heat and pressure are applied so that in most cases the material plasticizes and flows to fill the mould. Some compression moulding processes do not require flow, as in advanced composite prepreg moulding. In conventional compression moulding one of the mould halves (usually female) is loaded with a moulding charge sufficient to produce an accurate replica of the mould form. The charge usually covers 3&70'/0 of the female mould cavity surface. As soon as the male/upper cavity surface comes into contact with the charge it is heated very rapidly by conduction from all surfaces of the hot mould and the material begins to flow, taking only a few seconds to fill the enclosed cavity completely. Shortly after flow ends, heat from the mould cures the resin matrix, producing a solid part in one to three minutes. The temperature and pressure requirements may vary considerably depending upon the thermal and rheological properties of the polymer. For most compression moulding materials operating conditions are within the ranges 135-170°C and 7-20 M N m - 2 pressure respectively. A slight excess of material is usually placed in the mould to ensure it is completely filled. The excess polymer is squeezed out between the mating surfaces of the mould in a thin, easily removed film, known as flash. The compression moulding process is a fairly slow, labour intensive and, therefore, quite expensive method of producing parts. Vacuum bag moulding is a variant used for complicated shapes. 3.4.2

Injection moulding

The prime advantage of injection moulding is that it is a highly mechanized and controlled process capable of resulting in fast production rates

62

Short fibre-polymer composites

and low labour costs, and, as a consequence, low component cost and consistent operation. This in turn gives good quality products and a low reject rate. Injection moulding is free of several of the disadvantages that limit compression moulding; there is no need to weigh the charge, and material handling is much improved using a feed funnel containing ready mixed materials. The cavity is rapidly and identically filled each time without any pre-cure stage. After injection and setting of the polymer the screw rotates and moves backwards to prepare the next charge. Meanwhile the mould is opened and the moulded article removed. Rates of production are fast, not only because of the speed of operation of the machine and the mechanized method of dispensing the material to the cavities, but also because the process delivery system heats the material efficiently both by conduction of heat from the barrel and the development of heat due to friction first caused by the rotational motion of the screw. The melt temperature may be considerably higher than in compression moulding. As a result of this the cycle times can be very short indeed. Another benefit is the improved surface/visual quality of injection moulded polymer composites. Compression mouldings often have an uneven finish in surface gloss and colour; they may also contain traces of dirt picked up during the moulding operation. Injection moulding provides vastly improved surface and visual properties of mouldings by eliminating contamination and providing consistent and repeatable conditions appropriate to achieving best surface finish. 3.4.2.I Injection moulding process requirements A maximum injection pressure of 140 MPa is normally sufficient to mould polymeric materials efficiently. The moulding machine should have a non-return valve to prevent back flow of molten material. The barrel should be fitted with a water jacket so that heating may be applied. A maximum barrel temperature of 90 "C is adequate for most commercially available resins. Mould temperatures similar to those used for compression moulding, namely 135-170 "C,are used, the fast operation of the injection process allowing the use of high temperatures without the risk of pre-cure. The barrel may be maintained at ambient temperature; increasing the temperature to 60 "C allows faster injection speeds, some reduction of the cure time and better strength retention in the product. For the best surface finish, fast injection speeds are required and the fill time should be within the range 0.5-3.0 s. 3.4.2.2 Injectionlcompression moulding Automobile manufacturers and other engineering companies are currently considering increased use of thermoset composite materials because these materials can aid noise reduction, save weight and have a longer life.33 At one time thermosets were regarded as unacceptably difficult to mould and limited in their applications, but new formulations, new process techniques, fully automated equipment and advanced control systems have made these composites highly competitive. In the injection/compression procedure a charge of material is injected into a partially open mould (5-10mm) for the full injection stroke, then the mould is

Thermosetting short fibre reinforced composites

63

fully closed for the required curing time. The completion of mould closing after the injection step causes additional material flow within the mould, ensuring optimal filling of corners and intricate spaces. This process gives some additional control of fibre orientation. This is a result of the two stage moulding process. In the initial stage of the moulding procedure (the injection stage), some fibre orientation related to the cavity shape and the pressure can occur. During the compression stage the fibres reorient under a pressure that is applied equally across the whole section and high degrees of fibre orientation are generally not obtained. As a consequence the properties are not as anisotropic as with mouldings possessing highly developed fibre orientation. In some applications this may be more beneficial than the presence of the strong anisotropy found in components in which pronounced preferred orientation produces very good properties in one direction but significantly inferior properties in other directions. This process also minimizes the weld line size in multiply gated mouldings. Correct mould design is essential for the production of accurate specimen dimensions with injection/compression moulding. For injection/compression moulding the configurations are different from conventional tools in that the cavity is completely enclosed even at the parting line. This containment prevents material leakage and is achieved by telescoping a portion of the cavity side into a matching clearance on the core side of the mould. Precision fit of the telescoping areas creates the cavity pressure needed for the compression cycle. 3.4.3 Reaction injection moulding

The most exciting development in recent years in injection moulding has come from reaction injection moulding (RIM) and reinforced reaction injection moulding (RRIM).The processes are designed, as the names suggest, for reacting systems, particularly polyurethanes, but also to a more limited extent for other polymers such as epoxy resins and polyesters where no appreciable volatile products are formed during cure. Very fast cycle times in the region of a few minutes are possible with this process even when making large parts since the moulding operation step is combined with the chemical reaction step which is catalysed so that the reaction occurs quickly. The ingredients, including the filler, are melted into a mixing head. The magnitudes of viscosity, temperature and pressure in the mixing head are controlled to better than 1YO.Mixing is through turbulence and after each injection the mixing chamber is cleaned using a ram. Reinforced reaction injection moulding requires lower pressures than conventional injection moulding, leading to cheaper tooling costs. Another reason why RRIM is cheaper than injection moulding is because lower clamping pressures are required, reducing the cost of the moulding machine and making RRIM energetically favourable. 3.4.4 Hand laminoting

Hand laminating, or as it is also known ‘contact moulding’, is a traditional fabrication technique for GRP (glass reinforced plastic) materials. Chopped glass

64

Short fibrepolymer composites

rovings (short fibres) can be sprayed together with a controlled quantity of resin from a special gun on to the mould and then consolidated. The resulting 'spray up' material has a random distribution of fibres in the plane of the moulding, giving similar mechanical properties to chopped strand mat (CSM) polyester. Curing the resins can be speeded up by increasing the temperature of cure. It will take approximately one week to cure a moulding at room temperature but this can be reduced typically to 30 h at 40 "C or 3 h at 80 "C depending on the type of resin used. The main advantage of hand laminating is versatility. There is no size limitation. It is also a very adaptable production method since metal inserts or extra fibres can be added when they are needed. Because it is a low pressure, ambient temperature, process, inexpensive moulds of wood, plaster or GRP may be used which makes it suitable for small production runs. The main disadvantages are that mouldings usually have only one good surface and the process is labour intensive, with the quality dependent on the operator who controls the glass content, thickness and cure schedule, which can lead to variability in mechanical properties in a single moulding or between mouldings in a batch. The selection of suitable glass for hand lay reinforcement is carefully controlled. This is because most glasses on the market have been developed for polyester use. When a phenolic matrix is used, a different glass may be required. The sizes and binders used on these glasses do not affect the polyester chemistry but have some adverse effects on the phenolics. For reinforcement of phenolics by the hand lay-up process the reinforcement used is in the form of woven rovings.

3.5 Reinforcing fibres Fibre reinforced materials have been around for thousands of years.49 The earliest fibrous reinforcements were straw or horsehair for toughening mud in order to make bricks. Modern composites are reinforced with fibres such as glass, carbon or Kevlar. There are also particulate reinforcements using materials such as silica flour, glass beads and sand and there are now even lamellar reinforcements possible. Glass fibre reinforcement is commonest and is dealt with exclusively here. 3.5. I

Glass fibre reinforcement

The most common glass fibres for reinforcement are based on E-glass (E = electrical grade) which is a calcium alumina borosilicateglass. The fibres are made by heating the glass and drawing it through a platinum mandrel. The mandrel contains approximately 200 holes and so 200 fibres are made at each press. The hole diameter is accurately machined to give fibre diameters of 5-20pm. The glass is a good reinforcing medium because it has a high tensile strength (up to 3.6GNm-') and a high Young's modulus value (as high as 94GN m - 2 ) . Glass fibres have good dimensional stability and do not exhibit creep, and they bring these same benefits to composites that contain them. Since glass is a relatively inert material, the fibres will also be immune to biological

Thermosetting short fibre reinforced composites

65

attack and have good resistance to most forms of chemical and solvent attack. Glass fibres are also completely fireproof and have good weathering and electrical properties. Another advantage and possibly one of the most useful is that glass fibres are available in different lengths and diameters, and in different bundle sizes or as single fibres. Early attempts to use glass fibres to reinforce materials were disappointing owing to the brittle nature of the glass, causing breakage during moulding at high pressures and in injection moulding during processing in the screw/barrel. Steps to overcome this include passing the fibres through a sizing bath. The size provides protection to the fibres and also binds them into a strand, but does not bond adjacent strands together, thus allowing the strands to be easily unwound from the cake at later processing stages. The most common type ofsize used with thermosetting polymeric systems is based on poly(viny1 acetate). The incorporation of a resin coupling agent improves the bond between the laminating resin and the sized glass fibre and improves the mechanical and electrical properties of the reinforced material. By using glass fibres as a reinforcement, the strength and toughness of polymeric materials is vastly improved.

3.5.2 Behaviour offibres in resins During polymerization the exothermic cure reaction causes the temperature of the resin and therefore that of the glass fibres embedded in it to rise. When the reaction is complete the resin and fibres cool towards ambient temperature. As the thermal expansion coefficient of resin is approximately ten times that of the glass fibres, considerable compressive forces act on the glass on cooling, provided stress-free adhesion has occurred at the maximum cure temperature. It has been found that at a temperature characteristic for each resin, the actual compression forces are such that the fibre kinks, causing alternate compression and tension zones on the concave and convex sides of the fibre respectively. The kinking is reversible and disappears when the characteristic temperature is reached. High curing temperatures lead to high stresses and therefore an increased degree of kinking.50 Immediately following the curing procedure stress concentrations of both tensile and compressive nature exist along the glass fibre/resin bands. Fibre kinking is not usually desired in a composite because the initial tensile loads on the material will be taken entirely by the resin until the glass fibres are straightened. 3.5.3 Damage

to glass fibres during processing

The damage to glass fibres during the production of composite materials and the effect on the quality of the injection moulded parts has been the focus of considerable work.’ 1 9 5 2 The influence of the processing parameters and the injection moulding machine on the glass fibre distribution affects the mechanical characteristics of the injection mouldings. Glass fibre debris is increased at high injection velocities and with high dynamic pressures. Screws that have suffered

66

Short fibrepolymer composites

3.2 (a) Section through a reinforced phenolic injection moulding showing strong fibre orientation in the flow direction (perpendicular to the plane of the section). Iibres lying in this direction give circular images. The large circles correspond to glass spheres which are also present (fly ash). (b) Section through a DMC injection moulding in a region where the majority of the fibres are tilted away from the normal to the plane of the section, giving elliptical profiles.

excessive wear not only give a lower throughput and less reproducible settings but also increase fibre fracture, causing a reduction in mechanical properties.

3.6

Fibre orientation

Fibre orientation occurs with thermoset injection mouldings in much the same way as in thermoplastic injection mouldings (Fig. 3.2). There is a greater tendency

Thermosetting short fibre reinforced composites

67

3.3 Section through an injection moulding made from reinforced phenolic, showing fibre clumping and zoncs with very low fibre content.

for fibre clumping in thermoset mouldings (Fig. 3.3) with the consequence that sections used for fibre orientation analyses of the kind described in section 1.5.1 are less representative than is the case with fibre reinforced thermoplastics and the results show more scatter. Correspondingly, properties such as Young’s modulus, calculated from the orientation distribution, are less representative and less likely to agree with macroscopic measurements made on the mouldings. It is possible to have varying degrees of flow of fibres and/or of melt in compression moulding processes for thermoset reinforced moulding.53 In sheet moulding compound (SMC) the fibres are initially randomly oriented. However, as the mixture becomes fluid in the mould i t deforms and the deformation changes the orientation of the fibres. Orientation distributions can be extremely complex. Some locations can retain random orientation, whereas others may have highly aligned fibres. Advani and Tuckers4 have modelled the flow of fibres in a compression mould using a numerical simulation. In compression moulding,46 as soon a s the upper platten touches the charge, the material begins to flow and in only a few seconds the cavity is completely full. Fibres in the initial charge should be almost randomly oriented, then as the fibre/resin mixture flows, deformation occurs and this changes the orientation of the fibres. Shortly after flow ends, heat frotn the mould cures the resin producing a solid part in one to three minutes with fixed fibre orientation. Fibre orientation is influenced by the shape of the charge and the position in which it is placed, by the flow characteristics of the moulding compound, by the moulding conditions, and by the type of gating system used. Fan gates may give considerably more transverse fibre orientation than edge gates. The rate of filling through an edge gate is of no consequence to orientation distributions possible, but through a fan gate the fill rate can be important; filling slowly through a large fan gate or quickly through a small fan gate produces high longitudinal orientation. The

68

Short fibre-polymer composites

orientation becomes more longitudinal due to squirting, i.e. fast fill through small gates. In injection moulding additional factors may influence the fibre orientation in thermoset composites. Rojas et alS6studied the flow of thermosets through the nozzle of an injection moulding machine. They found that partial resin cure may occur; in extreme cases the gel point may be reached, causing partial blocking, disturbing the filling process, and ultimately preventing the operation. Rojas et alS6 attempted to model the process and deduced that most of the pressure drop occurs in the outlet section. With a non-reactive polymer they found that the pressure dropped because the temperature increased and the viscosity fell correspondingly. With a reactive polymer the effect of crosslinking is to increase the viscosity, opposing the effect caused by the increase in temperature, and the pressure drop is greater. Owen and Whybrew5’ found that, in both transfer and injection moulding, as the material enters the mould through the narrow gate, fibres become aligned in the flow direction before embarking on various expansions and contractions, dividing and rejoining as the mould fills. In regions i n which the mould section contracts, the fibres become oriented parallel to the flow direction, whereas in regions in which expansion occurs the fibres become aligned perpendicular to the flow direction. A more rigorous investigation of mould filling of injection moulded thermosets was attempted by Gunther and K r e t ~ c h m a r . ’ ~ Computerized methods have been used for designing moulds to injection mould reinforced thermoplastics (see section 1.3) but these procedures have not reached the same level of development with thermosetting moulding compounds. In experiments with short shots, thermoset composites exhibited different behaviour from thermoplastics processed in the same way. The flow fronts tended to break up, which may be the consequence of much lower melt extensibility in stretching flow. The reduced melt extensibility may be due to high filler content and/or the different shape of the melt front. A similar break up of the flow front can be found with highly filled thermoplastics but to a lesser extent. Gunther and K r e t ~ c h m a rmade ~ ~ several tests to compare thermoplastic and thermoset materials by moulding samples in a cavity containing a stepped region. They found that when thermoplastic was used the stepped region filled at the same time as the rest of the mould, but with thermosets there was a marked delay in the filling of the stepped region. They explained this difference in terms of the different flow processes taking place at the flow front with thermosets and thermoplastics: 1 Thermoplastics: because the mould temperature is less than the melt temperature, a cooled viscous outer layer develops next to the mould wall and constantly re-forms as the melt front advances. 2 Thermosets: because the m.ould temperature is higher than the melt temperature, the outer layer is heated and flows easily. This effect is enhanced by additional warming by the high shear developed in this area. The inner melt regions therefore slide on this outer layer. When the compound reaches the end of the channel the flow movement of the outer layer is stopped. If this is a

Thermosetting short fibre reinforced composites

69

step in a terraced mould, there is then an interruption offilling until a flowable outer layer is formed in the new flow direction. Because of these differences thermosets need higher pressures to push the material around the edge and to form a new flowable layer. For thermosets, overflow occurs from the side walls to the step region, whereas for thermoplastics overflow occurs from the step region onto the side walls. A change in mould temperature has practically no influence on mould filling. When injection moulding comparable types of materials with different viscosities and curing times, identical flow behaviour is recorded: Filler content does affect the flow pattern, however, and using lower filler contents causes the flow front to become more closed. Flow fronts also display a bulge in the side wall regions. Owen et studied the injection moulding of tensile test bars from DMC. Careful observation of these bars showed that the fibres were highly aligned with the flow in the gate but they rotated soon after passing through the gate to become aligned in circular arcs centred at the gate. Once the material touched the sides of the mould the flow became directed along the mould, forming stagnation zones at the corners together with an edge boundary layer. A restriction of flow in a mould acts as a gate. The dominant features of fibre orientation in thermoset composites on injection moulding are:

1 2 3 4 5

Strong alignment of fibres on flowing through the gate; Rapid realignment of fibres with expanding circular arcs on leaving the gate; Directed flow with shear planes and edge boundary layers; Fibres aligned parallel to the weld line when flow fronts meet; Surface layers with non-typical fibre orientation.

Uncured DMCs possess elastic properties, and material expanding freely from a gate can be regarded as having membrane properties due to the interlacing of fibres in the expanding front. Once the expanding front touches the sides of the mould, the development of the shape of the membrane can be envisaged to occur in one of two possible idealized ways, according to whether or not slip occurs at the boundary. In the slip state the membrane has infinite mobility. Segregated slurry initiates membrane slip. In the no slip option the membrane is pinned at one end of the cavity and as it stretches along the cavity walls, the portion between the walls has a radius centred on the gate. As the no slip membrane expands, material is pulled away and dumped at the sides to form a boundary layer at the edge and material is pulled away from the centre via membrane slip. With injection/compression moulding, fibre orientation effects due to mould shape can be nullified since in the injection part, flow of the material into the mould cavity is very important, but during the compression stage the fibres can be reoriented and in many cases tend to become random. Interactions between fibres can also influence orientation properties by tending to cause random fibre orientation to occur and so the volume fraction of fibres in the matrix can be quite critical. Injection moulding causes damage to the fibres and as a result the parts may

70

Short fibre-polymer composites

not perform as well as those made by compression moulding. Improvements to the injection moulding process are possible and Gibson found that by using a deep flighted screw to minimize work done on polyester-based composite DMC, fibre degradation during passage through the barrel was greatly reduced.60 ‘Low work screws’ are now widely available. It has been found that the initial fibre orientation before moulding strongly affects the orientation of fibres after moulding. The requirements for achieving acceptable moulded products from the injection moulding process are: (i) that the fibre orientation in the starting material should be random; and (ii) that the compound should have low viscosity under processing conditions. The latter requirement is achieved by using a polyester resin additive, permitting controlled manufacture ofacompound that is thick enough to aid handling but which under injection moulding conditions undergoes a considerable decrease in viscosity, so allowing it to be processed without undue fibre breakage. As with the incorporation of short fibres into thermoplastics, their use in thermosets generally improves their mechanical properties and the property changes depend on the fibre orientation distribution and the fibre-matrix interaction. Short fibre reinforced thermosets can be fabricated using the common processing techniques employed for unreinforced thermosets, but their final properties are usually more sensitive to the processing conditions than those of the corresponding unfilled polymers. The flow behaviour of the melt has the greatest influence on the distribution of fibre orientations in the final product. Studies by Goettler6’ on the fibre orientation distribution within reinforced epoxy mouldings and the dependence of orientation on the production method used showed that edge gates gave a preferred fibre orientation in the flow direction whereas fan gates gave a more transverse Orientation. The fill rate was found to have little effect on the fibre orientation in edge gated cavities, a fast fill rate through a small fan gate gave preferred fibre orientation in the flow direction. This was thought to be due to ‘squirting’, in which the molten polymer front does not flow smoothly within the mould cavity. The same fibre orientation in the flow direction was also observed with slow fill rates through a wide fa2 gate. Low viscosity melts were found to develop a higher degree of fibre orientation in the flow direction than higher viscosity melts.

3.7 Fillers and other additives to thermosets The OPEC oil crisis in the mid-1970s made the price of polymers increase enormously. There was much incentive to seek cheap fillers that could be added to polymer materials to reduce the volume cost without compromising the properties. Mineral fillers such as calcium carbonate, calcium sulphate, clay, aluminium silicate and alumina trihydrate were used to help decrease the amount of expensive polymer being used. The fillers tend to control the viscosity of the system and also reduce any tendency for the reinforcement to separate out during processing. Although the motive for including filler is often simply cost-saving, fillers can sometimes be incorporated to provide specific properties such as

Thermosetting short fibre reinforced composites

71

tracking resistance or fire retardance. The size and shape of filler particles can influence surface finish and also fibre degradation during processing.60*62-68 An important feature of thermosetting composite production is shrinkage control. Initially thermosetting materials had very poor shrinkage resistance which led to poor surface quality of mouldings, but the development of low profile additives has allowed moulding to much closer tolerances. During curing, unsaturated polyester resin tends to shrink away from the mould walls.64 The high polymerization shrinkage force creates internal stresses that may lead to warpage, sink marks, internal cracks, blisters, poor dimensional control and poor surface quality (often associated with fibre protrusion). To compensate for the resin shrinkage, and to eliminate most of the other problems, low profile additives are widely used. Several authors have proposed mechanisms to explain the action of the low The most likely source of shrinkage compensation is profile the formation of microvoids in the interfacial region between the resin and the additive phase (fibres or filler).66*67 This is believed to be most likely to occur when there is sufficient low profile additive to cause the separation of the matrix polymer and the thermoplastic low profile additive into co-continuous phases. When a tensile stress is applied to the co-continuous structure it cavitates relatively easily at low strains. The internal stresses generated by differential shrinkage in the vicinity of the fibre or filler particles cause this effect: experiments by Bucknall and co-workers showed that shrinkage compensation did not occur when filler was absent (and no other sources of forces were Traditionally poly(methylmethacrylate) (PMMA) and poly(viny1 acetate) (PVAC) have been used as low profile additives. The success of these low profile additives has resulted in the increased use of reinforced thermosetting materials, particularly polyester-based products. This class of material can now be used for automotive exterior panels where the surface must be cosmetically perfect and tough enough to stand the rigours of assembly and use. Other uses of polyester composites in the automotive industry include boots, bonnets, sunroofs, grills and headlamp housings. The use of low profile systems also gives other advantages to the moulding operation, including the possible incorporation of cores and inserts, leading to reduced costs. Lower moulding costs are possible if GRP moulds are used (instead of metal moulds) and lower assembly costs because complicated components can be moulded in a single piece. The disadvantages in low profile systems are the need for preforming, which is time consuming and hence costly, and the need for trimming, which generates problems in waste disposal and has an added time factor. Low profile additives also improve the toughness of the composite. The effect of low profile additives on residual stress is dealt with in section 3.8. In addition to polyester-based composite materials, phenolic resin composites can also be improved by the use oflow profile additives. When phenolic materials are reinforced they have good stiffness, strength, fatigue resistance, thermal expansion and environmental performance. Their main limitation is low impact

72

Short fibre-polymer composites

strength and toughness. By adding linear high molecular weight thermoplastics, which are inherently tough, the system can be significantly improved.

3.7.I

Coupling agents

Another useful addition to polymer composites is a coupling agent which provides a coating on the fibre that develops a strong bond between fibre and matrix. A brief introduction to the use of coupling agents in thermoplastic composites was given in Chapter 2 (section 2.2.4): the basic chemical action is the same in thermoset composites and is summarized below. The surface of glass consists of randomly distributed groups of oxides with a composition that depends on the composition of the glass. Similar groups are present on the surface of mineral fillers used in polymer composites. Some oxides, e.g. SiO,, Fe,O, and A1,0,, are hygroscopic and absorb water as hydroxyl groups and as molecular water, which is held to the hydroxyl groups by hydrogen bonding. Other oxides are hygroscopic and become hydrated. Thus glass picks up water very easily to form a well-bonded surface layer which may be many molecules thick. Long contact with water results in the solution of hygroscopic elements leaving a porous surface made up of a network of non-hydrated oxides. The formation of a water layer cannot be avoided in commercial processing and the coupling agent is usually applied in a water-based size. The presence of water leads to a large reduction in the surface energy and has an adverse effect on the wettability of glass fibres by the polymer matrix. Part of the function of the coupling agent is to increase the surface energy to ensure good wetting. The primary function of the coupling agent is to provide a strong chemical link between the oxide groups on the fibre surface and the polymer molecules of the resin. The general chemical formula for silicone-based coupling is R---SIX,. This is a multifunctional molecule which reacts at one end with the surface of the glass and at the other with the polymer phase. The group X bonded to the silicon represents hydrolysable groups bonded to silicon such as the ethoxy group, OC,H,. These are present only as intermediates since in the aqueous size solution they are hydrolysed to yield silanol R-Six, + 3H,O + R-%(OH), + 3HX. The trihydroxysilanols are able to compete with water at the glass surface by hydrogen bonding with hydroxyl groups at the surface (Fig. 3.4(a): M represents the metal element in the oxide involved in the coupling, Si, Fe, Al, etc). When the sized fibres are dried, water is removed and a reversible condensation reaction occurs between silanol and the surface between adjacent silanol molecules on the surface. The result is a polysiloxane layer bonded to the glass surface (Fig. 3.4(b)). Thus the silane coated fibre presents a surface of R-groups to the surrounding resin. In a properly designed system the organofunctional R-groups are chosen to be chemically compatible with the resin so that during the curing process they react with groups in the resin and become strongly bonded to it. The silane coating leads to a strong water-resistant bond.

Thermosetting short fibre reinforced composites

R

R

I

HO-Si-OH

I

H‘”

R

I

I

HO-Si-OH

.o

H’”

\.H

HO-Si-OH

I

.o

\.H

H”’

M

.o

\.H \0:. I

\0.= I

\0:. I

73

M

M

........................................... glass

(4

R

I

R

I

R

I

- O-Si-O-Si-O-Si-OI

0

I

M (b)

I I

I

0

0

M

M

I

/////////////////////Ill/ glass

3.4 Action of coupling agent: (a) hydrogen bonding between hydroxide Sonds in silanol (coupling agent) and metal hydroxide in glass (or mineral filler); (b) polysiloxane bonded to the glass (or mineral) surface.

3.8 Residual stresses Residual stresses in thermoplastic mouldings were discussed in section 2.5.5. Stresses of this kind develop when temperature gradients exist while a moulding sets, and in thermosets they are sometimes called curing stresses. The temperature gradients present when a thermoset cures are quite different from those in a thermoplastic moulding while it is solidifying. Schulze et a173 found that the temperature variations in thermoset mouldings have the effect that surface layers cure (crosslink) first, then form a rigid cage which prevents shrinkage, particularly in the injection direction. An additional compaction of moulding zones in the vicinity of the gate takes place when there is rapid injection with a melt cushion and high follow-up pressure. Using a finite element method to assess stress levels in the mouldings it was found that the largest relative stresses occur in the longitudinal edge of the moulding.73 The magnitude of the residual stresses depend on the moulding material, the mould temperature and the mould cavity pressure, and some degree of control is possible, especially if the moulding process has good cavity pressure monitoring and feedback control. Residual

Short fibre-polyrner composites

74

stresses may continue to change after the moulding has cooled, especially if it comes into contact with an environment with which the material interacts. The changes are unlikely to be as great as those observed with thermoplastics because relaxation in the heavily crosslinked thermoset is much more restricted.

3.8. I

Residual stresses in filled thermosets

At all locations within a thermoset moulding there will be a ‘microstress’ field around the reinforcing fibres or particles because of the differential thermal contraction of the filler and matrix. The shrinkage of the polymer is much greater than that of the filler and this exerts a stress normal to the surface of the filler which increases the frictional force against pull-out during fracture and thus enhances the toughness (see Chapter 1). There are in addition long range residual stresses that arise owing to the temperature gradients within the mould, as discussed above. They are of extreme importance with respect to design and performance of composites as they can lead to distortion of finished products and may assist the process of crack initiation and propagation. If the composite absorbs water the matrix swells and the thermally induced residual stresses decrease. In thermosetting matrix composites residual stress development is influenced by the process history.74 Processing effects such as curing, cooling rates and post-moulding operations are related to residual stress distributions over the thickness of a moulded sample.75These effects may be accentuated during the non-isothermal processing of polymeric materials causing solidification under a distribution of thermal and rheological conditions from the surface to the centre of the moulded sample. The most important contribution to residual stress formation in thermoset composites arises from the curing process, as found by Kau and P e t r ~ s h a Crasto , ~ ~ and Kim77 and White and Hahn.78,79The curing process determines not only the moulding cycle and therefore the production rate, but also the surface and mechanical quality of moulded parts. Residual stresses generated in the cure cycle are usually detrimental to mechanical strength and durability of thermoset composites. It is essential to control the thermal and mechanical histories of the material during the cure cycle and to analyse the chemical and physical properties of the material after it has been cured. Mallick and Raghupathi8’ found that the mechanical properties are affected by the moulding conditions, but these will affect other characteristics in addition to the residual stresses. There will be interaction between the degree of cure and the level of residual stresses and the variations in mechanical properties can rarely be attributed to one particular characteristic with confidence. 3.8.2

Measurement of residual stresses in thermoset composites

The authors have not found any reference to the measurement of residual stresses in short fibre reinforced thermosets in the literature and the following account is based on recent work conducted in their laboratory.

Thermosetting short fibre reinforced composites

75

Studies have been made of short fibre reinforced injection mouldings based on both a polyester resin and on a phenolic resin respectively. Residual stresses were measured by the layer removal procedure,81*s2using high speed milling to remove layers. This involves removing a thin layer from the surface of a flat moulding then measuring the curvature that results from the consequent imbalance in forces across the section as the moulding restores equilibrium of the internal bending moments, then repeating the process until at least half the moulding has been removed. A plot of curvature versus the depth of material removed is constructed and used to provide data from which can be calculated the distribution of residual stress in the original moulding, using equations provided by Treuting and In the case of unfilled polymers this curvature plot is reasonably smooth and leads to a fairly unambiguous derivation of the residual stress d i s t r i b ~ t i o n . ~When ~ - ’ ~ the same procedure was applied to fibre reinforced thermoplastic injection mouldings, the curvature plots sometimes showed more scatter, though it was still possible to locate a reasonably representative curve for a n a l y ~ i s . ~In ~ . the ~ ’ investigation of thermoset composites the curvature plots showed much more scatter, in some cases preventing the choice of a smooth representative curve for analysis. This led to critical examination of all aspects of the layer removal method and to minor improvements in the procedures. For example, a freshly ground tool was used for each layer removal, a precaution found to be unnecessary with thermoplastics. Even with the improvements, significant scatter still remained and it was necessary to devise data smoothing procedures to obtain curves suitable for analysis. The accuracy of this procedure is not easy to determine and the results provide no more than a guide to the behaviour; the numerical values cannot be regarded as accurate. The source of the scatter is believed to be the tendency for the fibres to clump, giving a strongly heterogeneous morphology. The fibre diameter was 8pm and a clump diameter is therefore a significant fraction of the sampling depth used in the layer removal tests ( - 100pm). Thus localized changes in property caused by the clumping are on the same scale as the measurement interval in the residual stress determination. A similar problem was encountered in the measurement of Young’s modulus distribution through the depth of the mouldings (section 3.9). The residual stresses measured in injection moulded reinforced phenolics generally showed distributions reminiscent of those found in thermoplastics, with a compressive region near to the surface and a broader tensile region of lower magnitude in the interior (Fig. 3.S(a)). To produce a stress distribution of this sense thermoelastically, the temperature should be highest in the interior at the completion of cure so that this is where the greatest shrinkage occurs when the moulding cools down to a uniform temperature. This would require exothermic heating to dominate heating by conduction from the mould. With DMCs the residual stresses tended to be tensile near the surface and compressive in the interior (Fig. 3.5(b)). In this case a temperature gradient in the opposite sense must be invoked to account for the stresses if they are formed thermoelastically. This implies that the curing exotherm does not reverse the temperature gradient

76

Short fibrepolymer composites

10

I

E

5

z I

.-

-10

a

Q, L

-15 0

1.o

0.5

depth

(4

1.5

(mm)

10

5

Y

0

-5 0 (b)

0.5

depth

1.o

(mm)

3.5 Residual stress distributions measured in injection mouldings made from (a) reinforced phenolic and (b) DMC. The analyses are shown for approximately half the moulding thickness (approximately 3 mm)."

1.5

Thermosetting short fibre reinforced composites

77

applied by the hot mould. An alternative explanation for the residual stresses could be formulated based on the shrinkage that occurs during the crosslinking reaction. If curing is effectively completed at different locations within the moulding at different times then the regions that cure last will contain tensile stresses; this will cause the regions which cured first to go into compression. To determine the origin of the residual stresses in thermosets would require knowledge of the temperature distribution throughout the moulding at all times during curing, and a detailed knowledge of the curing kinetics. It should be noted that a minority of phenoljc mouldings were found with tensile stresses at the surface and conversely a minority of DMC mouldings contained compressive stresses near to the surface. The stress magnitude was often quite a significant fraction of the breaking strength and it is evident that control over these stresses is required to ensure best mechanical properties. Lack of control may lead to unevenly balanced stresses across the moulding, and warping or distortion is fairly common. This can sometimes be remedied by placing the moulding into a jig immediately on ejecting it from the moulding machine, while it is still hot.

3.9

Properties of short fibre reinforced thermosets

Fibre reinforcement of thermosets provides at least as great property enhancement as with thermoplastics. Typical Young’s modulus values are in the range 10-1 8 GN m - ’. The properties of thermosetting composites are sensitive to the processing conditions, particularly in those based on phenolic resins, and can be highly variable even for nominally identical mouldings, and this limits their use.40*89-93 The present authors have investigated the properties of the mouldings at different depths and at different locations along the flowpath in injection moulded thermoset composites to attempt to explain some of the observations of property variation. 3.9. I

Young’s modulus distribution

The distribution of Young’s modulus through the thickness of the mouldings was determined by measuring the bending stiffness of the sample bar, then milling away thin layers from one side and repeating the measurement on the remainder. This process was repeated until about half the thickness had been removed. The measurement of bending stiffness was made either using a bending jig on a tensile test m a ~ h i n eor~ by ~ ,a ~resonance ~ method. A plot of the bending stiffness (e.g. the gradient of the force/displacement relationship in static bending) versus the depth of material removed was made and a third or fourth order polynomial fitted to it. This was used to determine the Young’s modulus distribution; the Young’s modulus at any depth is proportional to the gradient of this plot at the chosen position.88This procedure was important because of the scatter present in the measured data, believed to derive from the very heterogeneous nature of the

Short fibre-polymer composites

20

15

10

5

0 0

0.5

depth

1.o

1.5

(mm)

3.6 Young’s modulus distributions through the depth of two samples extracted from equivalent locations close to the gate in an injection moulded picture frame-type part made from a reinforced phenolic resin.88

mouldings. It permits the smoothing of data prior to processing and minimizes the errors that are normally associated with difference methods. In some earlier analyses, equal thicknesses were removed from both sides each time, and it was assumed that the Young’s modulus distribution was symmetrical about the centre plane of the m ~ u l d i n g . ~ ~ , ~ ~ The method involved plotting the bending stiffness of the bar remainder versus the depth of material removed. If the sample has a uniform or smoothly varying Young’s modulus this plot should be smooth, as was found with thermoplastics and fibre reinforced thermoplastic^.^^^^^ With the thermosetting composites the data points showed much more scatter but the data usually fitted a third or fourth order polynomial with a correlation coefficient better than 0.995. The corresponding Young’s modulus distributions are then derived and two examples are shown in Fig. 3.6. The results are for two different locations in the same reinforced phenolic moulding. The total depth of the moulding was about 3 mm and the distributions shown in Fig. 3.6 are for just less than half the depth. It is evident that there are large variations in Young’s modulus through the depth of the samples. The locations from which the samples were taken were equivalent to one another (equidistant from the gate, etc) and it is also clear that there are large

Thermosetting short fibre reinforced composites

79

20 N -

I

E 15

v)

10

5

0

depth

(mm)

3.7 Young's modulus distributions through the depth of two samples extracted from equivalent locations in injection moulded picture frame-type parts made from DMCs with slightly different compositions.88

differences in the distributions at different locations. Sometimes the largest value of Young's modulus was located near the surface and sometimes it was located in the interior. The method of measurement prevented accurate values from being obtained near the centre. Light optical study of sections showed that in many cases the fibres tended to be transversely oriented near to the centre and it is expected that Young's modulus will be smallest there. The results showed that there is a great deal of variability in the properties within phenolic injection mouldings that cannot be controlled by simple adjustment of the moulding conditions. Broadly similar results have been found with injection moulded DMC, Fig. 3.7. Young's modulus values at different depths in the same bar may differ by a factor of two or more, and even greater differences sometimes occur from one location to another in the same moulding. The variability almost obscured processing-dependent effects. Such large variations in property through the thickness will inevitably influence the overall engiiieering properties of the components. For example, the bending stiffness of two samples with the same average Young's modulus will be quite different if one has a strong variation in Young's modulus through the thickness and the other has a uniform modulus through the depth.

80

Short fibre-polymer composites

1750

I

"

"

I

'

I

I

I

1700 CI

m I

E cn

1650

&

1600

.-

1550

v

v)

c Q)

U

1500 1450

1400 0

0.5

depth

1.o

1.5

(mml

3.8 Density distribution through the depth of an injection moulding made from reinforced phenolic. The distribution is shown for half of the bar thickness.

3.9.2

Density measurement and cornposition investigation

Measurements of average density through the depth of the mouldings have been made on samples prepared by machining a rectangular parallelepiped from a moulding. The(average)density wasmeasured usinga relativedensity bottle then a layer was milled away from the surface and the density of the remainder was measured. This process was repeated until more than half the thickness had been removed. With large samples the average density was calculated from the weight and the linear dimensions instead of by density bottle. The density distribution through the depth was calculated from a plot of the average density of the remainder versus the total depth removed. It is a simple task to determine the composition of a composite containing two phases of known density from the overall density but quantitative interpretation is not possible when there are three or more phases present. Large variations in density were often observed through the depth of the bar (Fig. 3.8). In phenolic mouldings 3 mm thick it was found that the fibre content was often greatest about 1 mm from the surface. Partial segregation of the fibres also took place along the flow path, with the lowest concentrations being recorded near the gate and the highest concentrations furthest from the gate. With DMC mouldings

Thermosetting short fibre reinforced composites

81

the fibre concentration tended to be greatest near the centre but there was again a tendency for fibre concentration to increase progressively with distance from the gate.

3. I0 Conclusions Reinforced thermosets offer an extremely useful range of properties. They are stiffer than the most common reinforced thermoplastics and can be used at higher temperatures. Their disadvantages are that they are less convenient to mould and their properties are more variable and less easy to control. References 1 2 3 4 5 6 7 8

9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29

Ashby M F and Jones D R H, Engineering Materials I I , Pergamon, Oxford (1986). Piggott M R, Load Bearing Fibre Composites, Pergamon, Oxford (1980). Addabbo H E, Rojas A J and Williams R J J, Polym Eng Sci, 19 (1979) 835. Williams R J J, Rojas A J, Marciano J H, Ruzzo M M and Hack H G, Polym Plast Techno1 Eng, 24 (1985) 243. Castro J M and Straus E J, Polym Eng Sci, 29 (1989) 308. McGee S H, Polym Eng Sci, 22 (1982) 484. Kubota H, J Appl Polym Sci, 19 (1975) 2279. Lem K W and Han C D, J Appl Polym Sci, 28 (1983) 3185. Lem K W and Han C D, J Appl Polym Sci, 28 (1983) 3207. Lem K W and Han C D, Polym Eng Sci, 24 (1984) 175. Lee D S and Han C D, Polym Compos, 8 (1987) 133. Lee D S and Han C D, Polym Eng Sci, 27 (1987) 964. Joneja S K, Res Develop, 27(3) (1985) 84. Mostovoy S and Reppling E J, J Appl Polym Sci, 10 (1966) 1351. Kim S L, Polym Eng Sci, 18 (1978) 1093. Yamini S and Young RJ, J Muter Sci, 15 (1980) 1814. Parrat N J, Fibre Reinforced Materials Technology, Van Nostrand Reinhold, London (1972). Sheldon R P, Composite Polymeric Materials, Elsevier, New York (1982). Parkyn B, Glass Reinforced Plastics, Butterworth, London (1970). Holliday L, Composite Materials, Elsevier, London (1966). Thomas K and Dawson D, Pinst Rubber Proc Applics, 5 (1985) 293. Somerfield K T, Thermoset Injection High Performance at Low Cost, PRI Conference, 2-3 March (1982), Paper 16. Blanc R, Agassant J F and Vincent M, Polym Eng Sci, 32 (1992) 1440. Kelly A, Concise Encyclopaedia of Composite Materials, Pergamon, Oxford (1986) 113. Daspet Y, Thermosets 90’s: Meeting the Challenge, Regional Technical Conference, Rosemont, Illinois, SPE, Brookfield, Connecticut (1990) 133. Chen F and Jones F R, Reinforced Thermosets, Conference (1991) Paper 10. Gardizella A and Mueller R, Kunststojie, 7 (1990) 71. Fitts B and Rice B A, Thermosets 90’s: Meeting the Challenge, Regional Technical Conference, Rosemont, Illinois, SPE, Brookfield, Connecticut (1990) 143. Mattigetz R, How Concept Becomes Reality, 36th International Symposium, 136 (1991) 1916.

82

Short fibre-polymer composites

30 Forsdyke K L, Reinforced Plastics, 16th Congress BPF Reinforced Plastics Group, 16 (1988) 9. 31 Fenton B and Fitts B, High Technology Conferenceon High Temperature Polymersand Their Uses, Cleveland, Ohio (1989) 141. 32 Vynckier Trade Literature, Vynckier N V, Vyncolite Division, Nieuwevaart 51, B-9000 Gent, Belgium. 33 Peterson L, Polymer Composite Applications for Motor Vehicles, Society of Automotive Engineers Transactions: Technical Paper no 910048, SAE, Warrendale, PA (1991). 34 Bauschaus F and Benfer W, Adu Polym Technol, 2 (1982) 283. 35 Macosko C W and Gonzaleromero V M, Polym Eng Sci, 30 (1990) 142. 36 Atkins K E, Seats R L and Rex R C, Polymer Composite Applications for Motor Vehicles,Society of Automotive Engineers Transactions: Technical Paper no 9 10384, SAE, Warrendale, PA (1991). 37 Geelan V, Kunststoffe, 8 (1991) 258. 38 Gibson A G, Composites. 20 (1989) 57. 39 Bowyer G, PRI Conference, Thermoset Injection High Performance at Low Cost, March (1982) Paper 9. 40 Williamson A G and Gibson A G , Plast Rubber Proc Applics, 4 (1984) 203. 41 Goff J and Whelan T, Brit Plast Rubber, 40 (1986) 36. 42 Carrow S R, SPE 38th A N T E C (1980) 465. 43 Buschhaus F and Benfer W, Adu Polym Technol, 2 (1982) 283. 44 Rossi F, Bosco G, Rossetto E and Chiesi F,SPE 45th A N T E C (1987) 918. 45 Wehrenberg R H, Mater Engineer, 93 (1981) 46. 46 Eduljee R F, Gillespie J W Jnr and Pipes R B, Compos Sci Technol, 33 (1988) 241. 47 Parker F J, Prog Rubber Plasr Technol, 1 (1985) 22. 48 Parker F J , PRI Conference, Thermoset Injection High Performance at Low Cost, March (1982) Paper 8. 49 Ashby M F and Jones D R H, Engineering Materials 1, Pergamon, Oxford (1980). 50 Folkes M J and Russell D A M , Polymer, 21 (1980) 1252. 51 Obieglo G and Roller B, Kunststoffe, 76 (1986) 709. 52 Sandai A R and Piggott M R, J Mater Sci, 20 (1985) 431. 53 Lockett F J, Plast Rubber Proc Applics, 5 (1980) 85. 54 Advani S G and Tucker C L, Polym Compos, 11 (1990) 164. 55 Gibson A G , Thermosets 1991, Solihull, Birmingham, PRI, London (1991) Paper 9. 56 Rojas A J, Addabbo H E and Williams R J J, Polym Eng Sci, 21 (1981) 634. 57 Owen M J and Whybrew K, Plast Rubber, 1 (1976) 231. 58 Gunther P and Kretschmar 0, KunststofSe, 78 (1988) 677. 59 Owen M J, Thomas D H and Found M S, S P I 33rd Conference, Birmingham (1978) Paper 20B. 60 Moore D R, Carter J T and McGrail P T , Soc Automotiue Engs, Special Tech Publ, 812 (1990) 1. 61 Goettler LA, Modern Plast, 47 (1970) 140. 62 Stetzenberger H D, Compos Struct, 24 (1993) 219. 63 Tipping G, Conference, Reinforced Thermosets, Birmingham (1991) Paper 20. 64 Suspene L and Shaw Y Y, Plast Compounding, 14 (1991) 50. 65 Carter J T, Plast Rubber Proc Applics, 16 (1991) 157. 66 Bucknall C B, Partridge I K and Phillips M J, Polymer, 32 (1991) 636. 67 Phillips M J, Partridge I K and Bucknall C B, S A M P E Symposium on Advanced Materials: The Challenge of the Next Decade, April (1990) 21 11.

Thermosetting short fibre reinforced composites

68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95

83

Lam P W K, Polym Compos, 10 (1989) 439. Pattisson VA, Hindersinn R R and Schwartz W T, J Appl Polym Sci, 18 (1974) 2263. Pattisson V A, Hindersinn R R and Schwartz W T, J Appl Polym Sci, 19 (1975) 3045. Siegmann A, Narkis M and Kost J, Intern J Polym Muter, 6 (1978) 217. Demmler K and Lawonn H, Kunststoffe, 60 (1970) 954. Schulze V, Holweg S and Weise J, Kunststoffe, 83 (1993) 145. Bogetti T A and Gillespie J W Jnr, J Compos Muter, 26 (1992) 626. Manson J A E and Seferis J C, J Compos Muter, 26 (1992) 405. Kau H T and Petrusha LA, Polym Eng Sci, 30 (1990) 805. Crasto A S and Kim R Y, J Reinf Plust Compos, 12 (1993) 545. White S R and Hahn H T , J Compos Muter, 26 (1992) 2402. White S R and Hahn H T, J Compos Mater, 26 (1992) 2423. Mallick P K and Raghupathi N, Polym Eng Sci, 19 (1979) 774. White J R, Polym Testing 4 (1984) 165. Treuting R G and Read W T Jnr, J Appl Phys, 22 (1951) 130. Sandilands G J and White J R, Polymer, 21 (1980) 338. Haworth B and White J R, J Muter Sci, 16 (1981) 3263. Thompson M and White J R, Polym Eng Sci, 24 (1984) 227. Srivastava A K and White J R, J Appl Polym Sci, 29 (1984) 1241. Hindle C S, White J R, Dawson D and Thomas K, Polym Eng Sci, 32 (1992) 157. Wilkinson S B, P h D thesis, University of Newcastle upon Tyne (1996). Gibson A G, Plast Rubber Proc Applics, 3 (1983) 207. Johnson A F, Encyclopedia of Composite Materials, Pergamon, Oxford (1989). Chou T W and Kelly A, Annu Rev Muter Sci, 10 (1980) 229. Gibson A G and Payne D J, Composites, 20 (1987) 151. Thomas K and Dawson D, Plast Rubber Proc Applics, 5 (1985) 293. ODonnell B and White J R, J Appl Polyrn Sci, 47 (1993) 189. O’Donnell B and White J R, Plust Rubber Compos Procs Applics, 22 (1994) 69.

Short fi bre-t hermoplastic elastomer composites G B N A N D 0 A N D B R GUPTA

4. I

Introduction

With the advent of the polymeric age in the latter part of the present century, polymers and their composites are being widely used in sophisticated areas of application. Composites have gradually replaced traditional engineering materials such as wood, metals, glass and even ceramics in many applications because of their high strength to weight ratio, high design flexibility, ease of fabrication and relatively low cost. Because of these considerations, they occupy an enviable position in the polymer industry and have also come to the forefront of research and development activity in polymer engineering and technology. All plastics contain additives in one form or another, ranging from traces of catalyst residues to deliberately added mineral fillers in large proportion. The additives can be classed broadly as particulate or fibrous. The present chapter deals with the use of discontinuous fibres as fillers in thermoplastic elastomer matrices.' During the second half of the present century thermoplastic elastomers have emerged as an important class of material to substitute both for conventional thermoplastics and for elastomers in order to overcome the weaknesses inherent in these materials. The advantages of thermoplastics are ease and economy in processing, high strength and rigidity, and recyclability. Their disadvantages are higher set properties, lower flexibility and low thermal and low dimensional stability at elevated temperatures. On the other hand, elastomers have higher elasticity, flexibility, dimensional and thermal stability and resilience and lower set properties, but suffer from poor processability, high energy consumption, higher cost of production and non-recyclability. The thermoplastic elastomers act as a bridge between these two classes of polymers, combining some of the advantages of each category to give rise to a new class of materials. They combine the processing advantages of thermoplastics with the property advantages of elastomers. They are easily processed on standard thermoplastic processing equipment at elevated temperatures, and on cooling down to room temperature they behave like chemically crosslinked rubbers, without addition of any

Short fibre-thermoplastic elastomer composites

85

vulcanizing agents. Their other advantages include a wider range of hardness, excellent mechanical and rheological characteristics and transparency (leading to a wider choice of colours), and recyclability without sacrificing the original elastomeric properties. Therefore, the thermoplastic elastomers are economically much more attractive than individual rubbers and plastics. The typical elastomeric properties of thermoplastic elastomers (TPEs) are attributed to the physical crosslinking resulting from domain formation of triblock or multiblock copolymers having hard and soft segments. The hard segments remain incompatible in the rubber matrix, forming discrete microdomains and acting as physical crosslinks between the polymer chains at ambient temperature. At elevated temperatures, these thermolabile domains soften and the crosslinksdissociate,enabling the material to flow in a manner analogous to a thermoplastic. On cooling to room temperature the crystalline domains are reformed, thus acting as effective crosslinks and filler particles.

4.2

Classification of TPEs

TPEs have been broadly classified into two main categories:2 (1) block and graft copolymers and (ii) rubber-plastic blends. These are further discussed below. 4.2. I

Block and grafi copolymers

The graft copolymers are usually also block copolymers and the list of different species includes: 1 Polystyrene thermoplastic elastomers; 2 Thermoplastic polyurethane (TPU) elastomers; 3 Polyester thermoplastic elastomers; 4 Polyamide thermoplastic elastomers; 5 Graft copolymers. The polystyrene-elastomer block copolymers styrene-isoprene-styrene (S-I-S) and styrene-butadiene-styrene (S-B-S) were first introduced in 1965 by Shell Chemical Co, USA, under the trade name of Kraton. Thermoplastic polyurethane elastomers were introduced by Schollenberger et al of the B F Goodrich research laboratory in the 1 9 5 0 ~They . ~ consist of linear primary polymer chains comprising long flexiblechain segmentsjoined end to end by rigid chain segments through covalent chemical bonds. Polyols or macroglycols react with a diisocyanate in the presence of a low molecular mass glycol chain extender to produce linear block copolymers. The hard blocks have a glass transition temperature (T,) well above the normal ambient temperature, whereas T, for the soft blocks is much below room temperature. These are more closely related to the polyester or polyether thermoplastic elastomers. The polar nature of the recurring rigid urethane chain segments results in strong mutual attraction leading to extensive hydrogen bonding, aggregation and ordering of the crystalline domains leading to paracrystalline regions. In

86

Short fibre-polymer composites

addition there are weak van der Waals forces throughout the polymer chains which are long enough to entangle with each other. The polyester thermoplastic elastomers are segmented copolyether esters formed by transesterification of dimethyl terephthalate, a poly(alky1ene ether diol) such as poly(tetramethy1ene ether glycol) (PTMEG) (molecular weight 60&3000), and a low molecular mass diol such as 1: 4 butane diol. Long, hard segments of crystallizable tetramethylene terephthalate form and act as crosslinks that bind the soft amorphous elastomeric poly(alky1ene ether glycol terephthalate) into a network. Thus the whole system behaves like a conventional crosslinked elastomer. Polyamide thermoplastic elastomers consist of hard segments of adipic acid capped poly(I1-amino acid) and soft segments of poly(oxyethy1ene glycol) or poly(oxytetramethy1ene glycol). The hard segments have a crystalline melting point (T,) varying from 120 to 210 "C, and the flexible polyether segment a T, of - 60 "C. Some examples of graft copolymers are: polystyrene and poly(methy1 methacrylate) grafted natural rubber (NR) and butyl rubber grafted polyethylene.

4.2.2 Thermoplastic elastomers from rubber-plosric blends Thermoplastic elastomers are also prepared by melt blending a crystalline or semicrystalline thermoplastic with an amorphous rubber, giving rise to a two-phase morphology. The hard and solid phase is provided by the crystalline thermoplastic and the soft and flexible phase is provided by the amorphous rubber. The matrix is strengthened by the hard phase, hence the properties more or less depend on the nature and amount of the hard phase. This group covers a large number of thermoplastics and rubbers. In this class, tailor-made properties can be achieved for a particular end use by judicious selection of the polymers and by varying their ratios in the blends. This category of TPE includes miscible blends, immiscible blends and technologically compatible blends or alloys. Important among these are polyolefin thermoplastic elastomers, such as blends of NR-high density polyethylene (HDPE),NR-low density polyethylene (LDPE), NR-PP, ethylene-propylene diene modified (EPDM)-PE, EPDM-PP, LDPEpoly(dimethy1 siloxane) (PDMS), and other combinations include acrylonitrile butadiene copolymer-poly(viny1 chloride) NBR4PVC) and thermoplastic polyurethane (TPUkPDMS. Besides the above two categories, there are other types of thermoplastic elastomers, which can be termed as emerging thermoplastic elastomers. 1,2-polybutadiene with more than 90% vinyl content, ethylene vinyl acetate copolymer with 1&35% vinyl acetate content and ethylene ethylacrylate copolymer are also termed as thermoplastic elastomers. Similarly, thermoplastic elastomers in the form of natural rubber grafted with polystyrene or poly(methylmethacry1ate) are used commercially. Butyl rubber grafted on to PE is also an example of a thermoplastic elastomer. In general, thermoplastic elastomers are phase separated systems, where one phase influences the behav-

Short fibre-thermoplastic elastomer composites

4.1

87

Molecular structure of soft and hard phases in a thermoplastic elastomer. (Source: Ref 3, p 74.)

iour of the other. Therefore TPEs virtually bridge the property gap between thermoplastics and elastomers. The temperature range of application of TPEs is dependent upon the T, of the soft segment and T,,, of the hard segment, and the choice of the hard phase mostly decides the use of the TPE. The morphology of the hard and soft phases in a TPE is illustrated in Fig. 4.13and further discussion is given by Grady and C ~ o p e r . ~

4.3

Short fibre-elastomer composites 4.3. /

Short fibrerubber composites

Short fibre reinforcement in polymers has been used since the 1 9 6 0 Derringer6 ~~ incorporated different short fibres such as rayon, nylon and glass fibres into natural rubber matrix to improve Young’s modulus of the vulcanizates. He compared Young’s modulus of short fibre filled natural rubber vulcanizates with different types of fibre. Since then, short fibre reinforced rubber composites have gained considerable importance in the design of rubber products and short fibres have emerged as viable alternatives to particulate fillers in the reinforcement arena. Typical advantages associated with short fibres as fillers in polymer matrices include design flexibility, high low-strain modulus, anisotropy in technical properties and stiffness, good damping, ease in processing and production economy. Fibres may also improve thermomechanical properties of polymer matrices to suit specific areas of applications and to reduce the cost of the fabricated articles. Short fibre reinforced rubbers are generally easier to process than those with continuous cord reinforcement. The properties of articles made from short fibre reinforced elastomer composites depend on several factors such as the fibre aspect ratio (average length to diameter ratio), fibre orientation, fibre-rubber interface strength, state of dispersion, nature of matrix and type of fibre. The formulation of the rubber compound should facilitate processing and give good stress transfer.’-’ Moreover, short fibres provide high green strength, high dimensional stability during fabrication, improved creep resistance, good ageing resistance and damping, improved tear and impact strength and anisotropy in mechanical properties. The manufacture

88

Short fibrepolymer composites

of complex shaped engineering articles is impractical from elastomers reinforced with continuous fibres but is easily accomplished with short fibres. They are incorporated directly into the rubber compound along with other additives, and the compounds are easily processable in conventional rubber processing equipment, such as mixing mills, extruders, calenders and compression, injection or transfer moulding presses. Much work has been done on short fibre reinforced elastomer composites and the mechanism of reinforcement has been studied extensively.’ 3 - 1 6 Fibre length in a composite is a critical parameter. The term ‘short fibre’ means that the fibre length should not be too long, because the fibres may get entangled with each other, causing problems of dispersion, whereas a very small length of the fibre does not offer sufficient stress transfer area to achieve significant reinforcement. The term ‘composite’ signifies that the two main constituents, i.e. the short fibres and the rubber matrix, remain recognizable in the designed material. Goettler” has published a review on short fibre-rubber composites emphasizing the advantages and disadvantages of these materials. Ibarra and Chamorro’s.’9 reported the reinforcement of different elastomeric matrices with poly(ethy1ene terephthalate) (PET) fibres and found that the properties are dependent upon fibre concentration and type. 4.3.2

Short fibre-thermoplastic elastomer composites

The use of short fibre reinforcement in thermoplastic elastomers has not been very extensive hitherto but has recently opened up a new era in the field of polymer technology. It was first reviewed by Aoki” and by Theberge and Arkles.21 KaneZ2 studied the physicomechanical reinforcement of a butane diol-poly(tetramethy1ene glycol-terephthalic acid) thermoplastic elastomer by short glass fibres. Campbell and G ~ e t t l e r have * ~ reported the reinforcement of SantowebLbfibres in a TPE matrix. Akhtar eta124later reported reinforcement ofshort silk fibre in a T P E matrix based on blends of natural rubber and polyethylene. They studied the physico-mechanical properties and morphological features of these composites. Recently, reinforcement of short Kevlar.rhl aramide fibre in thermoplastic copolyester and thermoplastic polyurethane has been reported by Watson and Frances.2s Nangrani and Gerteisen26have also reported that the toughening of thermoplastic nylon 6,Ccarbon fibre composites could take place with the incorporation of a reactive thermoplastic elastomer and concluded that notched Izod, Charpy and Gardener impact strength could be significantly improved by the addition of the elastomer or by increasing the fibre length or by both in combination. Very recently Kutty et d2’ and Kutty and Nando28-34studied exhaustively the reinforcement of a thermoplastic polyurethane matrix by short Kevlar fibres. The processing and flow behaviour were studied as well as the mechanical, dynamic mechanical, thermal, adhesion and fire retardant properties of these

Short fibre-thermoplastic elastomer composites

89

composites. They also investigated the effect ofshort polyester (PET)fibres on the mechanical properties of thermoplastic p ~ l y u r e t h a n e . ~ ~ Subsequently Roy and c o - ~ o r k e r s ~studied ~ - ~ ~the rheological, hysteresis, mechanical and dynamic mechanical behaviour of short carbon fibre filled S-I-S block copolymers, and TPEs derived from NR and HDPE blends.

4.4 Parameters influencing the characteristics of short fibre-polymer composites

The reinforcement caused by short fibres in the polymer matrix is governed by the following parameters: (i) fibre dispersion, (ii) fibre-matrix adhesion, (iii) aspect ratio of the fibre, i.e. length to diameter ratio, (iv) fibre orientation and (v) fibre concentration. 4.4. I

Fibre dispersion

The primary requirement for obtaining a high performance composite is good dispersion of fibres in the matrix. The factors that affect fibre dispersion in polymer matrices are fibre-fibre interaction and fibre length. These factors also account for the tendency to agglomerate during mixing. For instance, naturally occurring cellulose fibres agglomerate during mixing due to hydrogen bonding. A pre-treatment of the fibres is therefore necessary to reduce the interaction between fibres and to increase interaction between the fibre and rubber. Such pre-treatments include making of pre-dispersions and formation of a soft film on the surface. Pre-dispersions of chopped fibres such as polyester, glass and rayon in a neoprene latex for better mixing into polychloroprene rubber or styrene butadiene rubber matrix have been successfully studied by Leo and J o h a n ~ s o n . ~ ~ Goettler4' reported that cellulose pulp may be dispersed directly into the concentrated rubber masterbatch or into a final compound if it is sufficiently wetted to reduce fibre to fibre hydrogen bonding. The fibres disperse rapidly during the early stages of the mixing cycle but the dispersion gradually slows down as mixing continues. Fibre dispersion is also improved with increased power input and mixing time. The fibre length should be small enough to facilitate good dispersion. According to Derringer,41 the commercially available fibres such as nylon, rayon, polyester and acrylic flock must be cut into smaller lengths, of approximately 0.4 mm, for better dispersion. Shen and Rains4' have stated that a dimensionless dispersion number N,, which is a function of rotor length, rotor diameter, rotor tip clearance, mixing chamber volume, rotor speed and mixing time, is a reliable scale-up parameter for short fibre mixing in polymers. 4.4.2

Fibre-matrix adhesion

Fibre to matrix adhesion plays a very prominent role in the reinforcement of short fibres in the polymer matrices. The fibre-matrix interface is important in

90

Short fibre-polymer composites

determining the mechanical, dynamic mechanical and rheological characteristics of thecomposites, since the stress transfer occurs a t the interface from the matrix to fibre. Though the mechanism of stress transfer is not clear, it has been postulated that it takes place through shearing at the fibre-matrix interface. In composites with low fibre-matrix adhesion, Derringer4' observed that a region of yielding occurs extending over a large portion of the strain range which is accompanied by low tensile strength and high permanent set. Recently Yano and c o - ~ o r k e r have s~~ studied the dynamic viscoelasticity of isoprene rubber with treated short cellulose fibre and have found that benzylation of cellulose fibre reduces the storage modulus (E') of polyisoprene rubber (IR) composite, whereas mercerization increases it. Kwon and c o - ~ o r k e r have s ~ ~ measured the fibre-matrix adhesion while studying the fatigue endurance and viscoelastic hysteresis properties of short fibre-rubber composites. Kutty and N a n d have ~ ~ reported ~ that chemically treated polyester cord-NR vulcanizates exhibited lower Goodrich heat build-up than the untreated PET cord-NR composites. Also NR matrix compounded with HRH (hexamethylene tetramine, resorcinal and hydrated silica) drybonding agent gave lower heat generation than even chemically treated fibre-rubber composites, owing to better interfacial adhesion between fibres and the matrix. Evidence of fibre-matrix adhesion has also been reported by Ibrarra and Chamarro," while studying the mechanical and dynamic properties of carbon and polyester fibre reinforced EPDM matrices. Akhtar et a124 observed that incorporation of dry bonding agents into the elastomer matrix improved mechanical properties of the composites due to an enhanced fibre-matrix adhesion. 4.4.3 Aspect ratio of the fibres

Whereas with continuous cord reinforcement the fibre ends have very little influence, in short fibre rubber composites they play a significant role in the determination of ultimate properties. The concept of a 'critical fibre length', over which the stress transfer allows the fibre to be stressed to its maximum, has been used to predict the strength of the composites. The mechanism of stress transfer between the matrix and fibres of uniform radius and length, has been analysed by Broutman and A ~ g a r w a lwho , ~ ~have given the following expression for critical fibre length, I,:

14.11 where d = fibre diameter, ofu= ultimate film strength and T~ = matrix yield stress in shear. It has also been emphazised that while comparing the various fibres of different diameters, aspect ratio is the main factor, not fibre length. The aspect ratio, i.e. the length to diameter ratio (LID)of the fibres, is a major parameter that controls the fibre dispersion, fibre-matrix adhesion and optimum

Short fibre-thermoplastic elastomer composites

91

performance of short fibre-polymer composites. If the aspect ratio of the fibre is lower than the critical aspect ratio, insufficient stress will be transferred and the reinforcement will be inefficient. Several a ~ t h o r s ~ ~have - ~ ’suggested that an aspect ratio in the range of 100-200 is essential for high performance fibre-rubber composites for good mechanical properties. Useful materials can be made with reinforcement fibres of much lower aspect ratio, however, and Chakraborty et alsl have observed that an aspect ratio of 40 gives the optimum reinforcement in the case of carboxylated nitrile (XNBR) rubber composite reinforced with short jute fibre. Murthy and De’2,52have reported that for jute filled rubbers, good reinforcement can be obtained with aspect ratios of 15 (natural rubber) and 32 (styrene butadiene rubber). Ibarra and Chamorro’* have shown that a carbonfibre EPDM system with reasonable properties had an LID in the range of 30-45. For synthetic fibres an aspect ratio of 100-500 is easily attained as they are available in the diameter range of 10-30 pm. Senapati et u153*54have shown that an aspect ratio of 227 gives good reinforcement with short polyester fibres in NR and an aspect ratio of 170 gives good mechanical strength in the case of short nylon 6 fibre reinforced NR vulcanizates. One should remember that the fibres should neither be too long to entangle with one another and cause dispersion problems nor too short to offer insuflicient stress transfer and give poor reinforcement. Quite often the fibres undergo size reduction during mixing on a mill or in a Banbury. O’Connor’ studied the extent of fibre breakage after both processing and vulcanization, and concluded that fibre breakage and distribution of fibre length occur in the uncured stock during processing and not during curing. The severity of fibre breakage depends primarily on two factors, namely (i) the type of fibre and its initial aspect ratio and (ii) the magnitude of the shear force generated during mixing. Glass fibres, being brittle, possess low bending strength and suffer severe damage during mixing, unlike cellulose and nylon fibres which are flexible and have a high resistance to bending. The shear force generated during mixing is particularly high in the cases where the compound viscosity is high and if this force exceeds the fibre strength then high fibre breakage results. De and co-workers9-’1 * 1 2 * 2 4h ave studied the degradation of reinforcing fibres in natural rubber, XNBR, NBR and thermoplastic rubber, and found that the breakage of jute and silk fibres is less than that of glass fibres. Senapati et u153.54have reported that fibre breakage with synthetic fibres like nylon 6 and PET during mixing into rubber on a mixing mill is almost negligible. The aspect ratio of nylon 6 fibres changed from 180 before mixing to 170 after mixing and that of PET fibres decreased from 251 before mixing on the mixing mill to 227 after mixing. Akhtar et ut24reported a moderate breakage of short jute fibres during mixing with NR-PE thermoplastic elastomer in a Brabender Plasticorder. Kutty and N a n d ~ ” .reported ~~ significant breakage of short Kevlar fibres during mixing in a Brabender Plasticorder in TPU matrix at 180°C and at 60rpm rotor speed for 9 min. The length of the fibres decreased from 6 mm to approximately 1.5 mm, the diameter remaining the same. The fibre length

92

Short fibrepolymer composites

v)

2

n .-

0.5

'c

'c

0

0.4

.-

0.3

CI

I -

1

0

2

'c

L Q,

0.2 -

n

5c

0.1 0

L

I

5.8

6.O

6.4

6.2

fibre length

6.6

6.8

(mm)

0.30 0.25

0.20 0.15 0.10 0.05 0

0

0.5

1.0

1.5

length

2.0

(mm)

2.5

3.0

93

Short fibre-thermoplastic elastomer composites

v,

Q, L

0.35

n .-rc 0.30 rc

0.25

: 0.10 E L

3

c

0.05

0

c

t1 0

0.5

1.0

1.5

2.0

2.5

3.0

4.2 Fibre length distribution of short Kevlar fibres: (a) before mixing; (b) after mixing for TPU containing lOphr (parts per hundred of rubber) of short Kevlar fibres; and (c) after mixing for TPU containing 40phr of short Kevlar fibres. (Source: Ref 55.)

distributions before and after mixing for TPU-short Kevlar fibre composites are shown in Fig. 4.2. They have also observed that the Kevlar fibres break down first and then tear apart during mixing due to high shear forces encountered in the Plasticorder and breakage follows a kinking mechanism. This is shown in Fig. 4.3. 4.4.4

Fibre orientation

The preferential orientation of fibres in the matrix is the key to the development of anisotropy in the matrix and application of the composites in various industrial products. For example, in V-belts the base compound has to withstand compressive forces while allowing sufficient flexibility in the axial direction. Therefore transversely oriented fibres are more suitable for this purpose. In composites in which the fibres are randomly oriented in planes parallel to the surface, swelling is restricted in both the length and the width directions and hence swelling takes place only in the thickness d i r e ~ t i o n . Therefore, ~ oil seals made out of them tighten after swelling. Short fibres get oriented preferentially in a particular direction when the

Short fibre-polymer composites

94

4.3

Scanning electron micrograph of a Kevlar fibre undergoing breakage through kinking mechanism during processing.

composite undergoes shear flow. The type of flow is determined by the processing techniques adopted, such as milling, extrusion and calendering. M ~ N a l l has y~~ reviewed in detail the orientation of short fibres in polymer matrices. Campbell57 reported that when the rubber matrix containing dispersed fibres is made to flow in a non-turbulent manner, the fibres are turned and aligned in the direction of the matrix flow. Boustany and Coran5’ reported that a high degree of fibre orientation could be achieved by repetitive folding and passing through a two roll mill. The effect of mill parameters such as number of passes, nip gap and mill roll speed on the fibre orientation was initially studied by Moghe who reported that nearly 60-70% fibres get oriented in the direction of the applied stress. Senapati et u153.54have observed that two passes of the short nylon 6 and PET fibre reinforced NR composites through the tight nip of a mixing mill is sufficient to orient most of the fibres in the mill direction. The effect of mill opening (nip gap) and friction ratio of the mill and temperature of the rolls on the orientation of short Kevlar fibres in T P U matrix has been studied by Kutty and N a n d ~ . ~It’ was observed that the lower the nip gap, the higher is the anisotropy in tensile properties of the composite, implying greater orientation of fibres. This has been presented in terms of the anisotropy index (ratio of property for compounds having fibres oriented in longitudinal direction to those oriented in transverse direction) which reduces gradually with increasing nip gap. During processing and subsequent fabrication of short fibre rubber composites, the fibres orient preferentially in a direction depending on the nature of the flow, e.g. convergent, divergent, shear or elongational, as explained by Goettler and c o - w ~ r k e r s , ~ ’If- ~the ~ flow is convergent, the fibres align themselves in the longitudinal direction, and if the flow is divergent, the fibres orient in the transverse direction. A detailed discussion on the design of extruder dies for

Short fibrethermoplastic elastomer composites

95

controlling fibre orientation is dealt with by Goettler et aiS9who have applied this technique in the manufacture of radiator hose from short cellulose fibre-EPDM rubber composites. During shear flow as experienced in a Brabender Plasticorder or a capillary rheometer, the fibre alignment may be random or unidirectional depending on the rate of shear. Kutty et d2' have shown that the fibres are uniformly distributed and entangled at lower shear rates, whereas at higher shear rates they are more oriented in the core region along the flow direction in a capillary rheometer than at the periphery. This has been observed for short Kevlar fibre filled TPU composites at 40 phr loading of fibres (phr, parts per hundred of rubber, is a descriptor of concentration, widely used in the rubber industry, whereby the relative amount of each ingredient in the mixture is expressed assuming the amount of elastomer(s) to be present to be 100 parts by weight). If the flow is elongational, then the fibre orientation takes place mainly in the direction of applied stress which is experienced during sheeting through the tight nip of a mill or during calendering. Calendering can achieve the same level of fibre orientation in the machine direction as conventional extrusion and this technique is being used in profile extrusion of tyre component strips, such as filler strips e t ~ . ~ ~

4.4.5 Fibre concentration

The concentration of fibres in the matrix plays a crucial role in determining the mechanical strength properties of fibre reinforced polymer composites. A lower concentration of fibres gives lower mechanical strength. This has been observed not only in rubbersj3 but also in T P E m a t r i ~ e s . ~ ~ . ~ e*behaviour , ~ ~ , ~ ~has T hbeen attributed to basically two factors: (i) dilution of the matrix, which has a significant effect at low fibre loadings, and (ii) reinforcement of the matrix by the fibres which becomes of increasing importance as fibre volume fraction increases. At low fibre content, the matrix is not restrained by enough fibres and highly localized strains occur in the matrix at low strain levels, causing the bond between fibre and rubber to break, leaving the matrix diluted by non-reinforcing, de-bonded fibres. At high fibre concentrations, the matrix is sufficiently restrained and the stress is more evenly distributed, thus the reinforcement effect outweighs the dilution effect.' 9*53s4,64 As the concentration of fibres is increased to a higher level, the tensile properties gradually improve to give a strength higher than that of the matrix. The concentration of fibres beyond which the properties of the composite improve above the original matrix strength value is known as the optimum fibre concentration. In order to achieve improvement in mechanical strength with short fibres, the matrix must be loaded beyond this volume fraction of fibres. This optimum fibre concentration is quite often found to lie between 20 and 30phr. This has been observed by Kutty and Nando28for short Kevlar fibre filled thermoplastic polyurethane composites, as shown in Table 4.1, and by Senapati et a/ for short nylon and PET fibre reinforced natural rubber

Short fibre-polymer composites

96

Table 4.1. Mechanical properties of short Kevlar aramide fibre-TPU composites

Mix

Orientation

Tensile strength (MPa)

TKO

L T L T L T L T L T L T

36.5 36.1 16.7 13.8 26.0 15.5 31.9 19.6 44.0 21.4 58.4 26.6

TK,o TK20 TK30 TK40 TK,,

Modulus at 20% (MW

Elongation at break (%)

Tear strength kNm-'

1.2 1.I 11.6 5.8 23.0 8.3 34.0 11.0 44.0 14.3 54.5 21.7

580 630 1.09 540 39 120 27 57 20 55 22 32

106 106 84 86 73 72 90 I1 96 75

Impact strength Jm-')

Hardness Shore A

-

81

-

-

88 2.15 1.15 I .48 1.16 1.38 1.43 1.36

-

(x

-

-

90 -

92 -

93 -

93

TK stands for thermoplastic polyurethane short Kevlar aramide fibre composites with the subscripts denoting the phr of fibres. L = longitudinal and T (Source: John Wiley and Sons, 1991, Ref 28.)

= transverse.

composite^.^^.^^ Quite often, at very high fibre concentrations, for instance at 40phr or beyond of short nylon 6 or PET fibres in N R matrix, the strength again decreases, because there is insufficient matrix material to adhere the fibres toget her. 4.5

Short fibrethermoplastic elastomer composites (SF-TPE)

4.5.I

Processing

The mixing of short fibres into TPEs is comparatively easy, but controlling the conditions of mixing is quite tricky and tedious, since the fibres are processed like thermoplastics in plastics processing equipment. Good fibre dispersion is the ultimate objective of any mixing process. Depending on the type of fibre, mixing may be either distributive or dispersive. Distributive mixing increases the randomness of spatial distribution of the minor constituents within the major base without reducing the size of the fibre, whereas dispersive mixing serves to reduce the agglomerate size. However, both phenomena occur simultaneously. Brittle fibres such as glass, jute, Kevlar, carbon and boron break down severely during mixing on a mixing mill and in an intensive mixer. Therefore, these fibres need more distributive mixing, whereas organic fibres such as cellulose, nylon, polyester and silk require more dispersive mixing due to their tendency to agglomerate during the process. Conventional mixers such as mixing mills and Banburys are used for mixing short fibres with rubbers, whereas intensive mixers like extruders, plasticorders, injection moulding machines and kneaders are used for mixing short fibres in thermoplastics or thermoplastic elastomers.

97

Short fibrethermoplastic elastomer composites

Table 4.2. EfFect of nip gap of the mill during sheeting of short Kevlar fibre-TPU composites on the tensile properties (TK2&

Property Tensile strength (MPa) Elongation at break (X) Tear strength (kN/m)

Fibre orientation

La T L T L T

Nip opening (mm) 0.45

0.63

0.91

1.21

1.48

32

31 14

28 15

(2.4)

(2.0)

27 16 (1.7)

19 12

16 71 73 63

28 15 (1.9) 19

80

79 75

13 (2.5)b 12 105 77 65

15 66

74 51

18 54

L = longitudinal and T = transverse orientation of fibres. Values in parentheses are anisotropy index values; defined as a ratio of the tensile strength in the longitudinal direction to that in the transverse direction. (Source: Elsevier Science Publishers Ltd, 1993, Ref 29.)

a

The first and foremost criterion for mixing short fibres into a thermoplastic elastomer matrix is to form a homogeneous dispersion in order to achieve ultimate strength and performance of the composites. The second criterion is to effect a preferential orientation of the short fibres in the polymer matrix, to achieve anisotropy in properties. These are made possible by manipulating the mixing technique and controlling mixing parameters. The mixing process is optimized such that the breakage of the short fibres during mixing is minimal and the orientation of fibres is maximized. The effect of nip gap on the fibre orientation and hence on the anisotropy of tensile properties of TPU-Kevlar fibre composites has been studied by Kutty and N a n d ~ . Table ’ ~ 4.2 shows that on increasing the nip gap, the tensile strength in the longitudinal direction of the composites steadily decreases, whereas that in the transverse direction increases. As a result, the anisotropy index also decreases. The tear strength exhibits a similar trend. This has been attributed to the lower extent of fibre orientation as the nip gap is increased. The effect of friction ratio on the orientation of fibres in short Kevlar fibre filled TPU composites has been studied by Kutty and N a n d ~ . ’The ~ effect of friction ratio on the tensile and tear properties are presented in Table 4.3. An increase in friction ratio during final sheeting first increases the strength then decreases it. At a low friction ratio (1.05: 1) the shear force is not suficient to orient the fibres in the direction ofapplied stress, and at relatively high friction ratio (1.75: l), there is slip on the roll surface, as a result of which fibres are not sufficiently oriented along the mill direction. At the optimum friction ratio of 1.15: 1 the tensile strength as well as tear strength values exhibit the maximum anisotropy, indicating optimum orientation of the fibres in the T P U matrix under the conditions studied. A fairly modest effect of roll temperature on the properties of Kevlar reinforced thermoplastic polyurethane was observed by Kutty and Nando2’ who used mill roll temperatures in the range 30-100°C. In the case of short silk fibre filled thermoplastic elastomers derived from a blend of NR and

98

Short fibre-polymer composites

Table 4.3. EHect of friction ratio of the sheeting rolls during sheeting of short Kevlar fibre-thermoplastic polyurethane composites on the tensile properties (TK,,)

Property Tensile strength (MPa) Elongation at break (YO) Tear strength (kNn1-I)

Friction ratio

Fibre orientation

1.05

1.15

1.30

1.45

1.60

1.75

L T L T L T

30 1s 29 70 75 80

32 14 13 105 78 70

30 15 26 102 76 78

29 15 26 112 64 96

29 15 28 94 69 64

26 15 26 100 61 59

L = longitudinal and T = transverse orientation of fibres. (Source: Elsevier Science Publishers Ltd, 1993, Ref 29.) Tcible4.4. Mixingsequenceand fibredispersionofshort silk andcarbon fibresintoa blend of NR and HDPE Mixing sequence

I 11 111

IV

Dispersion of fibres

+ NR + short fibre NR + short fibre + H D P E

Molten H D P E

Not good; fibres are not uniformly incorporatcd Lumpy agglomerate of fibres and non-uniform dispersion

Molten H D P E + half fibre t NR + another half of the fibre Molten H D P E

Better dispersion than I and I I , few lumps formed Good, uniform dispersion of fibres

+ short fibre + rubber

(Source: John Wiley and Sons, Inc, 1986, Ref 24.)

HDPE, final sheeting was done on a two-roll mill at 85 f 5 "C to ensure optimum orientation of the short fibres in the grain d i r e ~ t i o n . 'Mixing ~ sequences for short silk fibre reinforced NR-HDPE matrix are given in Table 4.4. Sequence IV gave the best physico-mechanical properties. The same trend has also been observed by Roy et in mixing short carbon fibres in NR-HDPE thermoplastic elastomer. Kutty and Nando28 have provided a typical mixing cycle for short Kevlar fibre-thermoplastic polyurethane composites, carried out in a Brabender Plasticorder at 180 "C and at 60 rpm rotor speed. They found that adding the fibres to 50% of the molten mass and then adding the rest of the polymer gave the best dispersion. The effect of processing on fibre length distribution was discussed in section 4.4.3. 4.5.2

Rheology

A knowledge of the rheological behaviour ofshort fibre filled polymer composites is of paramount importance in understanding the behaviour of the composites during processing and forming operations. The literature contains a large

Short fibre-thermoplastic elastomer composites

2.0

1.6

I

I

E

I

I

I

o 160'C

I

I

*

,

1

180'C

0

99

'

1

1

200'C

> .............&.................................... ............ y,,, = 401s-' I

0

2

,

4

6

8

10

volume fraction of fibres

12

14

(%)

Variation of viscosity with fibre loading at direrent temperatures and rates of shear for short carbon fibre filled thermoplastic elastomer derived from N R + HDPE blends. (Source: Ref 38.)

4.4

number of references on the flow behaviour of short fibre filled thermoplastics as well as elastomers65 7 0 but there is little information about the rheology of short fibre filled thermoplastic elastomers. Kutty et dZ7 investigated the rheological behaviour of short Kevlar fibre filled thermoplastic polyurethane composites. They reported the variation of rheological characteristics with fibre concentration, temperature and rate of shear. Roy and c o - w o r k e r ~ ~have ~~~ reported ' the rheological behaviour of short carbon fibre filled composites of (i) thermoplastic elastomeric blends of natural rubber and high density polyethylene (NR-HDPE) and (ii) thermoplastic elastomers based on styrene-isoprene-styrene block copolymer (S-I-S). All the aforesaid composites show pseudoplastic or shear thinning rheological behaviour. The shear viscosity has been reported to increase with fibre loading for the short Kevlar fibre-thermoplastic polyurethane composites at low shear rates,27 whereas both the carbon fibre filled systems show different behaviour. At a low rate of shear, the carbon fibre NR-HDPE filled system38 gave only a marginal increase in viscosity up to 10 phr fibre loading, beyond which it either remained constant or decreased. However at a higher rate of shear, the viscosity did not show any increase with fibre concentration but remained practically constant, showing a decreasing trend at higher fibre loadings (Fig. 4.4). The carbon fibre filled S-I-S composite, on the other hand,

I00

Short fibre-polymer composites

gave a reduction in shear viscosity3’ up to about 10phr fibre loading beyond which it increased with increase in the fibre content. This behaviour is a consequence of two opposing factors, i.e. (i) an increase in viscosity due to the filler-matrix and filler-filler interactions and (ii) a decrease in viscosity due to wall slip caused by the alignment of fibres in the direction of flow. The overall effect of these two factors is different for different systems, thereby accounting for the variations in the nature of the viscosity change with the fibre loading. For short fibre filled NR-HDPE and S-I-S composites, the variation in viscosity has been correlated with filler loading, rate of shear and temperature by the following equation: 7 - 3 8

+,

whereq, and q p are the viscosities of filled and gum compounds, is the volume fraction of the fibre in the composites, and K , and K , are constants, functions of the apparent shear rate and temperature. The activation energy of the flow for short Kevlar fibre-polyurethane composites has been reported by Kutty et alZ7to decrease with increasing fibre loading and shear rate. At low shear rates the reduction in activation energy is more pronounced than at high shear rates. This implies better processability of the composites at higher fibre loading. The die swell characteristics of thermoplastic elastomer-short Kevlar fibre filled composites have been reported by Kutty et al,27 and short carbon fibre filled TPE systems by Roy and c o - w ~ r k e r sThe . ~ ~die ~ ~swell ~ value (ratio of extrudate diameter to the die diameter) has been reported to decrease with increasing fibre content. This is because of reduction in the amount of polymer matrix per unit volume of the compound, which gives a corresponding reduction in stored elastic energy with increase in the fibre content, and also because of polymer-fibre interaction.

4.5.3 Mechanical properties 4.5.3.I Tensile properties The mechanical properties of short fibre thermoplastic elastomer composites are intermediate between those containing continuous cords and those with particulate fillers. The theory of short fibre reinforcement was dealt with in Chapter 1. The variation of tensile strength with loading of fibres varies from rubber to rubber (and more so from thermoplastic to thermoplastic). For strain crystallizing rubbers, such as NR and CR, the tensilestrength drops initially up to a certain volume fraction of fibres, typically about 10 phr, even if there is good bonding, then it increases gradually with fibre l ~ a d i n g .For ~ ~non-crystallizing ,~~ rubbers, such as SBR and NBR, the presence of a small quantity of fibre enhances the strength.” Thermoplastic elastomers of block copolymer type, whose matrix strength is

101

Short fibre-thermoplastic elastomer composites

50

e

40

-

TK 40 (L)

/4 30 -

/

I / I

, I

20

10

0

strain 4.5

(%I

Stress-strain behaviour of short Kevlar fibre-thermoplastic polyurethane composites at room temperature (25 "C).(Source: Ref 28.)

much higher owing to strong intermolecular interaction between thechains in the hard phase, exhibit similar behaviour as strain crystallizing rubbers. Thermoplastic polyurethanes have a higher tensile strength (37-40 MPa) due to strong hydrogen bonding between urethane segments of adjacent chains in the hard phase, which act as crosslinks between the soft, flexible and rubbery polyol segments. On incorporation of short Kevlar fibre in proportions of 10,15 and 20 parts by weight, the tensile strength reduces drastically to 17, 24 and 26MPa respectively in the longitudinal direction2*This has been explained as due to dilution of the matrix and introduction of flaws at the fibre ends where high stress concentration occurs, causing bonds between fibre and rubber to break. At higher loading of Kevlar fibres, i.e. beyond 20 phr, the short fibres reinforce the matrix, facilitating easy transfer of stress which outweighs the dilution effect. The stress-strain behaviour of short Kevlar fibre-TPU composites is shown in Fig. 4.5. The elastomeric behaviour of neat TPU and the thermoplastic behaviour of the composite is depicted in this figure. The tensile strength in the transverse direction is always lower than that in the longitudinal direction as shown in Fig. 4.6 and 4.7. Similar observations have also been made by Akhtar et a124 for NR-LDPE and NR-HDPE blends filled with short silk fibres. Kutty and Nando3' studied the effect of temperature on the stress-strain behaviour of short Kevlar fibre filled-TPU composites. Figure 4.8 gives the

I02

Short fibrepolymer composites

" V

.El

40 -

400

In

0

E a .-

P

e.

300

30

In

C

s

20

a

c L .

T

In

P

x

elongation at break

100

10

-

.....,.... , . . . . . . . . . . . . . . .......

0 7

C

0

-bQ

In

., .--

200

20% tensile modulus

0, C

2

2

tensile strength

Q)

c.

lD 2

0

I

I

10

20

40

30

..

50

4 4

fibre content

(phr)

4.6 Tensile properties of short Kevlar fibre-TPU composites versus fibre loading with the fibres oriented in the longitudinal direction. (Source: Ref 28.)

I

SO+..,

In

-

3 3 -0

0

E a -

1

'. ..

40

I

"(.

-

tensile strength

v..,

ca

, 30 -',

600

I

I

0

20% tensile modulus

T

elongation at break

-

500 3

In

-

0)

400 1:

-'.,

0 3

1

300

2 rT (d

200 P

X

2 .- I n 0 c 0

0

10

20

0

40

30

50

CI

fibre content

(phrl

4.7 Tensile properties of short Kevlar fibre-TPU composites versus fibre loading with the fibres oriented in the transverse direction. (Source: Ref 28.)

Short fibre-thermoplastic elastomer composites

I03

stress-strain behaviour of the neat TPU (TK,) and T P U filled with 20phr of short Kevlar aramide fibres (TK,,) at 25, 50, 75 and 100°C, respectively. The modulus and tensile strength are reduced as the temperature is increased, indicating a gradual softening of the matrix. Elongation at break is increased for the unfilled TPU (TK,) when the temperature is increased, whereas the elongation at break is not very sensitive to temperature in TK,, (containing 20 phr of short Kevlar fibres). A comparison of short Kevlar fibre filled T P U composites with carbon black filled T P U composites7’ shows that incorporation of high abrasion furnace (HAF)black into the TPU matrix at loadings varying from 10 to 40 phr decreases the tensile strength of the composites at all black loadings. This has been attributed to the fact that black particles act as stress raisers rather than reinforcing filler in the TPU matrix. The tensile strength of T P U reduces from 37 to 18 MPa on incorporation of 40 phr of HAF black, whereas in the case of 40 phr Kevlar fibre filled TPU composites the tensile strength increases from 37 to 44MPa in the fibre orientation direction and decreases to 21 MPa in the transverse orientation showing a clear anisotropy in strength.” The modulus at 100% elongation of HAF black filled TPU composites is very low, 7 MPa at lOphr black loading and still only 12 MPa when the black loading is 40phr, in contrast to short Kevlar fibre reinforced TPU composites which have a relatively high modulus in the fibre orientation direction at 10 phr fibre loading of 11 MPa, and 44 MPa at 40 phr fibre loading even at 20% elongation. In the transverse orientation the moduli at 20% elongation are nearly 50% less than the former. 4.5.3.2 Tear strength In general, the tear strength of the short fibre reinforced polymer composites should improve to a greater extent when the fibres are oriented perpendicular to the direction of propagation of tear in the polymer matrix than when oriented parallel to it. This has been demonstrated by several workers for short nylon 6 and PET fibre filled natural and synthetic rubber ~ o m p o s i t e s .Sc ~ anning ~~~~,~~ electron microscopy (SEM) studies of tear fracture surfaces support this view. The tear strength is not considerably influenced when the tear propagates parallel to the fibre alignment direction, because the short fibres do not significantly contribute to the obstruction of the growing tear. However, in the case of short Kevlar fibre reinforced millable polyurethane rubber vulcanizates, tear strength improves significantly in the transverse ~ r i e n t a t i o n . ~ ~ If a thermoplastic elastomer matrix is used instead, as, for instance, in the case of short Kevlar fibre filled T P U composite^,^^ the tear strength reduces with increase in fibre loading up to 20 phr of fibres both parallel and perpendicular to the fibre orientation direction, then it almost stabilizes. The tear strength of neat TPU is found to be the highest (106 kN/m). This has been attributed to a large amount of stress concentration at the apex of the angular specimen producing triaxial stress. Thus application of uniaxial stress initiates the failure at the apex and propagates across the sample width. Since stress dissipation near the tip of the growing crack by viscoelastic processes is essential to the development of high

I 04

Short fibre-polymer composites

30

25

Q

B

20

Y

15 v) v)

al

L CI

ffY

10

5 0

25

15 v) v)

0, v)

10

5 0 0

10

20

30

4.8 (a) Stress-strain behaviour of TPU at elevated temperature. (b) Stress-strain behaviour of TPU filled with 20phr of Kevlar fibres at elevated temperatures.

Short fibrethermoplastic elastomer composites

I05

strength, the presence of the hard segments in the matrix results in an additional mechanism by which strain energy is dissipated. In addition to this, dispersed particles may serve to deflect or arrest the growing cracks, thus further delaying the failure. A comparison of tear strength of short Kevlar fibre filled TPU composites with that of the HAF carbon black filled TPU composites at identical concentrations shows that the tear strength of the latter composites is much higher (by 50%) than that of the former in the longitudinal orientation of fibres7' This increase in tear strength of the carbon black filled TPU composites as compared to the fibre filled composites is due to higher dissipation of tear energy and local deformation at the crack tip resulting in strain-induced crystallization and to crack blunting.73 4.5.3.3

Impact strength The impact strength of short fibre thermoplastic elastomer composites is reduced with fibre incorporation both parallel and perpendicular to the fibre orientation. The impact strength of TPU composites filled with PET or Kevlar fibres reduced drastically from about 25 J m- without short fibres to a value in the range 1-2 J m- with short fibres in proportions varying from 10 to 40 phr.35*71 This may be explained as due to higher low strain modulus of the composite; fracture initiates at the weak fibre-matrix interface before significant viscoelastic energy dissipation takes place.

'

4.5.3.4 Hysteresis and set

Short fibre filled thermoplastic elastomers show anisotropy in hysteresis loss as well as tension set. Carbon fibre filled thermoplastic elastomers based on NR-HDPE blends exhibit high levels of anisotropy in both these properties36 whereas carbon fibre filled S-I-S systems have very low anisotropy in these properties. The anisotropy in the hysteresis loss for a carbon fibre filled NR-HDPE system has been found to depend on fibre loading and strain. The reasons for the occurrence of anisotropy in the properties has been explained on the basis of level of randomness in the orientation of the fibres. It has been e ~ t a b l i s h e dthat ~ ~ the fibre orientation in filled S-I-S systems is more highly random than in NR-HDPE thermoplastic elastomer composites. 4.5.4

Dynamic mechanical behaviour

A knowledge of the dynamic mechanical properties of short fibre reinforced polymer composites over a wide temperature range is especially important for applications such as V-belts, conveyer belts and tyre components. Earlier studies have revealed that the dynamic modulus of the composites increases with short fibre incorporation and the loss peak broadens while showing a steady decline in rnagnit~de.~~-*~*~" The dynamic mechanical properties of short Kevlar fibre thermoplastic polyurethane composites has been studied by Kutty and NandoZswho varied the concentration of fibres. The storage modulus (E') in the fibre orientation direction

I06

Short fibrepolymer composites

was found to increase with fibre loading and this effect was more pronounced in recently observed the post-glass transition temperature (> T J region. Roy et that short carbon fibre filled TPE composites derived from blends of NR and HDPE exhibit anisotropy in the dynamic modulus similar to that observed for the static modulus of the composites. Also it has been observed that composites with transversely oriented fibres (10 phr carbon fibre loading) exhibit a lower modulus than those with longitudinally oriented fibres. The effect of fibre loading on the storage modulus and the loss tangent (tan 6) of short Kevlar fibre filled TPU composites in the longitudinal direction is shown in Fig. 4.9. The tan a,, value of the composites gradually decreases with increase in fibre loading, whereas the glass transition temperature (T,) remains unaltered. The transition peak progressively broadens with fibre loading which has been ascribed to polymer-filler interaction. The polymer in the immediate vicinity of a filler particle can be considered to be in a different state from that at remote locations. This may affect the relaxation of the matrix, resulting in a broader loss tangent peak.75This has been confirmed from the study of the loss tangent peaks of composites which do not show any change in shape where there is no interaction between the filler and the matrix, as reported by Dutta and T r i ~ a t h y for ’ ~ glass bead filled elastomeric composites. A similar trend has been observed by Roy et a136for short carbon fibre filled TPE from NR-HDPE blends where no change in shape (or position) of the transition peak occurs. T h i s suggests that short carbon fibres have virtually no interaction with the blend matrix. This had been earlier observed by Medalia77and later by Carling and Williams78for short fibre filled rubber vulcanizates. In this case tan6 is higher in the transverse orientation, much nearer to that for the unfilled blend than in the longitudinal orientation. This has been attributed to mechanical anisotropy. This has been observed for all loadings of short carbon fibre filled NR-HDPE blends at all frequencies.

4.5.5 Stress relaxation behoviour Stress relaxation is an important property in rubber technology for the rubber vulcanizates are subjected to various stresses in service, often under a wide range of temperatures. This has a significant influence on the service life of products which are subjected to static stress (hose, cables) as well as those which are subjected to dynamic stress (V-belts and tyres). Stress relaxation of NBR composites, filled with short jute fibre has been studied by Bhagawan et ul.” Subsequently FlinkBoconducted a detailed study on the stress relaxation behaviour of NR vulcanizates filled with short cellulose 4.9 (a) Storage modulus (E‘)versus temperature of short Kevlar fibre-TPU composites with variation in fibre loading with fibres oriented in the longitudinal direction. (Source: Ref 28.) (b) Effect of fibre loading on the loss tangent ofshort Kevlar fibre filled TPU composites with fibres oriented in the longitudinal direction. (Source: Ref 28.)

I07

Short fibre-thermoplastic elastomer composites

1O ’ O

1 o9

1o8

10’

1o6

-

(‘C)

temperature 1

lo-’

1 o-2 -100

0

temperature

100

(‘Cl

200

I08

Short fibre-polymer composites

fibre and grafted cellulose short fibre by adopting a procedure initially developed by Tobolsky" which assumes that stress relaxation can be modelled by an array of relaxing elements with different characteristic times. He also investigated the use of a continuous relaxation spectrum, and observed a two step relaxation phenomenon. Kutty and N a n d have ~ ~ ~studied the stress relaxation behaviour of T P U composites filled with short Kevlar fibre with respect to fibre loading, fibre orientation, strain rate and level, and relaxation temperature. They presented plots ofo(t)/o(O)versus log t, whereo(t) and o(0) are the stresses at times t and zero respectively. Neat T P U showed two straight lines of unequal slope intersecting at a point, unlike that observed for natural rubber gum vulcanizate.81 The occurrence of two straight lines indicates that two relaxation mechanisms operate, one at shorter times (< 200 s) and another at a later stage.82 This has been observed at three different strain levels (lo%, 50% and 500Y0)in the case of short Kevlar fibre-TPU composites. The rate of relaxation is found to be independent ofstrain levels up to 100% extension for shorter relaxation times but at higher extension (500%) the rate of relaxation is increased by 65%. Similar observations have been reported earlier for rubber vulcanizates.83~84 This has been attributed to the progressive build-up of the crystalline phase in the stretched state. The second slope also increases with strain level and the crossover point shifts towards shorter times at higher strain levels, implying that the initial mechanism of relaxation is quickly exhausted under high deformation.

4.5.5.I Effect of fibre loading The short Kevlar fibre-TPU composites showed a three-step relaxation process. In the presence of 10-30 phr of fibres the initial rate of relaxation was reduced by 60%. Further additions of short fibres did not change the rate of relaxation significantly. However, at 40phr of fibres, the rate of relaxation showed a 100% increase over that at 20 phr in the composite. The first phase of relaxation has been attributed to orientation at the fibre-matrix interface. A similar observation has been made earlier by Bhagawan et The second slope was found to be higher in the case of fibre filled T P U composites and the third slope increased with fibre loading, indicating that the mechanisms that operate initially as well as during the later stages of relaxation are functions of fibre loadings. As the fibre content increases, the intermediate phase at the fibre-matrix interface also increases, which may give rise to a higher initial rate of relaxation. However, in the case of neat TPU, only two distinct modes of relaxation are observed. T P U consists of hard segments that act as 'virtual physical crosslinks' in the soft matrix in which they are dispersed. Thus the initial relaxation in the composite may be arising out of the combined effect of the orientation at the hard-soft and fibre-matrix interfaces, whereas the second phase of relaxation may be due to the flow of soft matrix under tension at longer times. At low fibre loading, even though the fibre matrix interface content is high, the matrix gets diluted, whereas at higher fibre loading, the matrix is restrained sufficiently to give a net reinforcement effect, which results in higher relaxation temperature. Similar

Short fibre-thermoplastic elastomer composites

I09

observations have been reported by Flink" for short cellulose fibre reinforced natural rubber vulcanizates. The rate of relaxation at longer times was found to increase with fibre loading. This has been explained by Mullins and Tobin" for a composite's overall extension as due to local deformation, which may be of the order of the breaking elongation of the 'soft regions' in the composite. The short fibres restrain the soft regions, thereby reducing the amount of soft regions available for deformation and resulting in a higher relaxation temperature. The third slope was also found to increase with fibre loading but remained less than that of the second slope at all fibre loadings. The first transition shifted to shorter times on incorporation of the short fibres and remained more or less independent of further fibre loading. The second transition was shifted to longer times with fibre loading. 4.5.5.2 € f e a of temperature Stress-relaxation studies were carried out at room temperature (25 "C)and three other temperatures (50,75 and 100 "C) with neat TPU and TPU filled with 40 phr short Kevlar fibres. In the case of the neat TPU it was observed that the initial slope decreased first with temperature, in agreement with the observation reported earlier.79r83However, at 100"C, the rate increased slightly. The second slope increased with temperature, and at a higher temperature a third slope became apparent. In the case of 40 phr short Kevlar fibre-TPU composite the short and long term relaxation rates increased marginally with temperature up to 50 "C and shot up at higher temperatures. The relaxation at longer times is predominantly due to chemical relaxation which is a function of temperature, causing the second slope to increase with temperature. The rate of the third mode of relaxation increased only marginally with temperature. The activation energy calculated for the second mode of relaxation from the plots oflog(s1ope)versus inverse temperature, was 8.5 kJ mol- for unfilled TPU and 3 kJ mol- for T P U with 40 phr Kevlar fibres,

'

'

4.5.6

Self-adhesion behoviour

The self-adhesion behaviour of short Kevlar fibre-TPU composites has been studied by Kutty and N a n d ~ Using . ~ ~ a substrate of neat TPU, the peel force obtained with samples containing 20 phr Kevlar fibres (TK,,) was much greater than withothercompounds(O,10,40phr)(Fig. 4.10). Thevaluesobtainedforneat TPU and for TK,, were similar, and that for TK,, was intermediate between these and that for TK,,. A similar result was obtained for peeling from substrates of the same compound, i.e. TK, on TK,. These measurements were made while keeping the contact temperature, contact pressure and contact time constant at 140 "C, 1 MPa and 15 min respectively. In other experiments it was found that as the contact time was increased from 3 min to 15 min, the peel adhesion strength increased abruptly from 93 to about 225 N.

I10

Short fibre-polymer composites

250

-

200

-

150 100 -

50 0

I

0

10

20

30

fibre content in one phase only 4.10

40

(phr)

Effect of fibre loading on the joint strength of the composites when the fibres are present in one phase only. (Source: Ref 31.)

This behaviour has been attributed to the interdiffusion of matrix, leading to migration of a small proportion of the fibres to the interface along with the matrix which helps in bridging the two phases at a lower proportion of fibres. As the fibre concentration increases to 20phr, quite a good number of short Kevlar fibres interdiffuse and migrate to the interface, virtually forming bridges between the T P U and TPU-fibre matrix resulting in good anchorage. These fibre bridges formed at the interface obstruct the growing tear during peeling, thus improving the peel adhesion strength of the joints. It is observed that the interdiffusion of fibres along with the matrix at the interface is facilitated by the reduction of the size of the fibres from 6 mm to 1-1.5 mm which has taken place during mixing in a Brabender Plasticorder. This in turn results in an increase in modulus at the interface. Therefore, the combined strength is enhanced and with the TK,,/TK,, specimens adhesion was so good that failure at thejoint was found to be cohesive, not interfacial. However, on increasing the fibre content further up to 40 phr, the peel adhesion strength decreases again because the presence of more short fibres at the interface dilutes the matrix and, therefore, failure occurs easily because of lower fibre to fibre adhesion and less interdiffusion of the matrix.

Short fibre-thermoplastic elastomer composites

4.11

Ill

LO1 test samples of T P U and T P U filled with short Kevlar fibres before and after testing. (Source: Ref 32.)

4.5.7

Flame retordoncy

It has been established that smoke and toxic gases (STGs) in a fire are the principal hazards and the main cause of casualties. STGs obscure visibility, and toxic materials from the combustion products cause congestion in humans, leading to death. Therefore the recent trend is to develop ways to reduce the generation of STGs in a fire and to improve the flame-retardant property of the materials used rather than combating fire itself. Recently, Kutty and N a n d investigated ~ ~ ~ the flammability of short Kevlar fibre-TPU composites with respect to fibre loading and various flame retardant additives. They reported that a drastic reduction in the generation of smoke occurs in the presence of short Kevlar fibres but there also occurs a marginal decrease in the limiting oxygen index (LOI) on incorporating short Kevlar fibres into the TPU matrix. The decrease in LO1 has been explained as follows: during the LO1 test, the neat T P U sample is kept vertical, and molten material drips down, depleting the burning head of the fuel. In the presence of fibres this dripping is arrested completely even at very low fibre loading (10 phr) as shown in Fig. 4.11. Thus, in the absence of dripping, the primary degradation products of TPU, such as polyols and isocyanates, are available for further burning, and as a result a lower LO1 is obtained. Smoke generation is drastically reduced in the presence of Kevlar fibres due to

I12

Short fibrepolymer composites

40 0 residue a t 500'C 0 residue aa tt 800'C 800'C 0 residue

-calculated

30

0

residue

-

1

a4

g

20

U .-

v)

Q) L

10

0

0

10

20

fibre content

30

40

(phr)

4.12 Char residues of short Kevlar fibre filled TPU composites a t 500°C and 800 "C respectively versus fibre loading. (Source: Ref 87.)

the formation of less volatile products during combustion, leading to high char formation. Thermogravimetric analysis of I 0 4 0 phr Kevlar fibre-TPU composites show higher char residues than the neat T P U at 800°C (Fig. 4.12). Kevlar fibres are found to be more effective than other flame retardants both with respect to the LO1 and the smoke density. Interestingly, Kutty and Nando3' have observed that the flame spread resistance of the T P U matrix is greatly improved by the presence of short fibres. Incorporation of short Kevlar fibres in the T P U matrix significantly increased the flame propagation time even at low fibre content (10 phr). Beyond this loading of fibres the flame spread time is improved marginally, i.e. the flame propagation is reduced. At 1Ophr loading the time to propagate the flame through lOmm down a vertical sample is increased from 5 t o 36s (longitudinal fibre orientation) and 44 s (transverse). Because of dripping, T P U experiences faster flame spread, but in the presence of short Kevlar fibres, dripping has been totally prevented and the samples undergo charring. It is also observed that at all fibre loadings, the resistance to flame spread is higher when the fibres are oriented across the sample length than when they are oriented along it. 4.5.8

Thermogra vimetric analysis

The sensitivity of thermoplastic elastomers to high temperatures, at which they

Short fibre-thermoplastic elastomer composites

I13

melt and flow, can be reduced by incorporating high heat and flame-resistant fibres. As Kevlar fibres have high thermal stability, they are expected to improve the thermal stability ofTPUs and this has been observed.86As the fibre content in the composite is increased from 10 to 40 phr, the onset of thermal degradation is shifted from 245 to 255°C. The residual weights after degradation of the TPU-short Kevlar fibre composites increased monotonically with concentration of short Kevlar fibres as shown in Fig. 4.12. The order of the degradation reaction of the TPU-short Kevlar fibre composites has been calculated by the Freeman and Carroll method and is found to be first order at all fibre loadings.

References 1 Sheldon R P, Composite Polymeric Materials, Applied Science Publishers, London (1992). 2 De S K and Bhowmick A K, Thermoplastic Elastomers from Rubber Plastic Blends, Ellis Horwood Series in Polymer Science and Technology, Chichester, U K (1990). 3 Walker B M, Ed, Handbook of Thermoplastic Elastomers, Van Nostrand Reinhold, London (1979). 4 Grady B P and Cooper S L, Chapter 13 in Science and Technology of Rubber, 2nd Ed, Eds J E Mark, B Erman and F R Eirich, Academic Press, London (1994) 601. 5 Anon, Technical Report Nos 31 and 34, Rubber Chemical Division, Monsanto Co, Louvianla, Neuve, Belgium. 6 Derringer G C, Rubber World, 165 (1971) 45. 7 Moghe S R, Rubber Chem Technol, 49 (1076) 1160. 8 O’Connor J E, Rubber Chem Technol, 50 (1977) 945. 9 Murthy V M and De S K, Rubber Chem Technol, 55 (1982) 287. 10 Goettler L A and Shen KS, Rubber Chem Technol, 56 (1983) 619. 11 Setua D K and De S K, J Mater Sci, 19 (1984) 983. 12 Murthy V M and De S K, J Appl Polym Sci, 29 (1984) 1355. 13 Dyzura E A, Intern J Polymeric Mater, 8 (1980) 165. 14 Rosen B W, Fibre Composite Materials, American Society for Metals, Metals Park, Ohio (1965). 15 Beatty J R and Hamed P, Elasfomerics, llO(8) (1978) 22. 16 Boustany K and Coran A Y, US Patent 3,696,364 (1970). 17 Goettler L A, in Handbook ofElastomers- New Developmentsand Technology, Eds A K Bhowmick and H L Stephens, Marcel Dekker, New York (1988). 18 Ibarra L and Chamorro C, J Appl Polym Sci, 37 (1989) 1197. 19 Ibarra L and Chamorro C, J Appl Polym Sci, 43 (1991) 1805. 20 Aoki J, Raba Daijesuto, 28 (1976) 13. 21 Theberg J E and Arkles B, Machine Des, 48 (1976) 113. 22 Kane R P, J Elastom Plast, 9 (1977) 416. 23 Campbell J M and Goettler LA, in Proceedings of PRI National Corzfererice on Short-Fibre Reiniorced Thermoplastics, Brunel University, Uxbridge, Middlesex, UK (1985) Paper 14. 24 Akhtar S, De P P and De S K, J Appl P d y m Sci, 32 (1986) 5123. 25 Watson K R and Frances A, Paper presented at the 132nd Meeting, ACS Rubber Division, Cleveland, Ohio, October (1987).

I14

Short fibre-polymer composites

26 Nangrani K J and Gerteison S R, in 42nd Annual Conference and Expo'87; Proceedings, Cincinnati, Ohio, February (1987) 5. 27 Kutty S K N, De P P and Nando G B, Plast Rubber Compos Proc Applic, 15 (1991) 23. 28 Kutty S K N and Nando G B , J Appl Polym Sci, 43 (1991) 1913. 29 Kutty S K N and Nando G B, Plast Rubber Compos Proc Applic, 19 (1993) 105. 30 Kutty S K N and Nando G B , J Appl Polym Sci, 46 (1992) 471. 31 Kutty S K N and Nando G B , J Adh Sci Technol, 7 (1993) 105. 32 Kutty S K N and Nando G B, J Fire Sci, 11 (1993) 66. 33 Kutty S K N and Nando G B, P l a t Rubber Proc Applic, 14 (1990) 109. 34 Kutty S K N and Nando G B , J Appl Polym Sci , 4 2 (1991) 1835. 35 Kutty S K N and Nando G B, Intern J Polymeric Muter, 19 (1993) 63. 36 Roy D, Bhowmick A K and De S K , Polym Eng Sci, 32 (1992) 971. 37 Roy D and Gupta B R, J Appl Polym Sci, 49 (1993) 1475. 38 Roy D, Bhattacharya A K and Gupta B R, J Elastom Plast, 25 (1993) 46. 39 Leo T J and Johansson A M, US Patent 4,263,184 (to Wyrough and Loser Inc) 21 April (1981). 40 Goettler LA, US Patent 4,248,743, 3 Feb (1981). 41 Derringer D C, J Elastom Plast, 3 (1971) 230. 42 Shen K S and Rains R K, Rubber Chem Technol, 52 (1979) 764. 43 Yano S, Hirose S, Hatakeyama H, Westerlind Band Rugadohl M, J Appl Polym Sci, 40 (1990) 657. 44 Kwon Y D, Beringer C W, Feldstein M A and Prevorsek D 0,Rubber World, 29 (1990). 45 Kutty S K N and Nando G B , Kautschuk Gummi Kunststoffe, 43 (1990) 189. 46 Broutman L J and Aggarwal B D, in Analysis and Perjbrmance of Fiber Composites, SPE/John Wiley and Sons, New York (1980). 47 O'Connor J E, Rubber Chem Technol, 47 (1974) 396. 48 Boustany K and Arnold R L, J Elastom Plast, 8 (1976) 160. 49 Coran AY, Boustany K and Hamed P, Rubber Chem Technol, 47 (1974) 396. 50 Boustany K and Coran A Y, US Patent 3,397,364 (to Monsanto Co) 10 Oct (1972). 51 Chakraborty S K, Setua D K and De S K, Rubber Chem Technol, 55 (1982) 1286. 52 Murthy V M and De S K, J Appl Polym Sci, 27 (1982) 461 1. 53 Senapati A K, Nando G B and Pradhan B, Intern J Polymeric Mater, 12 (1988) 73. 54 Senapati A K, Kutty S K N, Pradhan B and Nando G B, Intern J Polymeric Muter, 12 (1989) 203. 55 Li P C , Goettler L A and Hamed P, J Elastom Plast, 10 (1978) 59. 56 McNally D L, Polym Plast Technol Eng, 8 (1977) 101. 57 Campbell J M, Prog Rubber Technol, 41 (1978) 43. 58 Moghe S R, Rubber Chem Technol, 49 (1976) 1160. 59 Goettler LA, Leib R I and Lambright A J, Rubber Chem Technol, 52 (1979) 838. 60 Goettler LA, Lambright A J, Leib R I and DiMauro P J, Rubber Chem Technol, 54 (1981) 273. 61 Goettler L A and Lambright A J, US Patent 4,057,610 (to Monsanto Co) 8 Nov (1977). 62 Goettler L A and Lambright A J, US Patent 4,059,591 (to Monsanto Co) 1 Nov (1977). 63 Goettler LA, Sezna J A and DiMauro P J, Paper presented at ACS Rubber Div Meeting, Cleveland, Ohio, October (1981); Rubber World, 187 (1982) 33. 64 Rueda L I and Chamorro AC, J Appl Polym Sci, 37 (1989) 1197. 65 Crowson R J, Folkes M J and Bright P F, Polym Eng Sci, 20 (1980) 925. 66 Crowson R J and Folkes M J, Polym Eny Sci, 20 (1980). 67 Gupta A K and Purwar S N, J Appl Polym Sci, 30 (1985) 1777.

Short fibrethermoplastic elastomer composites

68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86

I15

Fujiyama M and Kawasaki Y, J Appl Polym Sci, 42 (1991) 481. Wang K J and Lee LJ, J A p p l Polym Sci, 33 (1987) 431. Murthy V M, Gupta B R and De S K, PIast Rubber Proc Applic, 5 (1985) 307. Kutty S K N, P h D Thesis, IIT, Kharagpur (1991). Kutty S K N and Nando GB, Intern J Polymeric Muter, 17 (1992) 235. Lee D J and Donovan J A, Rubber Chem Technol, 60 (1987) 910. Kamal M R, Song L and Singh P, Polym Compos, 7 (1986) 323. Rodok J R M and Tai C L, J Appl Polym Sci, 6 (1962) 518. Dutta N K and Tripathy D K, Kautschuk Gummi Kunststoffe, 42 (1989) 665. Medalia A L, Rubber Chem Technol, 51 (1978) 437. Carling M J and Williams J G, Polym Compos, 11 (1990) 307. Bhagawan S S, Tripathy D K and De S K, J Appl Polym Sci, 33 (1987) 1623. Flink P, P h D thesis, Dept of Polymer Technology, the Royal Institute of Technology, Stockholm, Sweden (1989). Tobolsky A V, Properties and Structure of Polymers, New York, John Wiley and Sons (1960). MacKenzie C I and Scalan J, Polymer, 25 (1984) 559. Gent A N, Rubber Chem Technol, 36 (1963) 397, 697. Voet A, Sircar A K and Cook F R, Rubber Chem Technol, 44 (1971) 185. Mullins L and Tobin N R, Proceedings tf the 3rd Rubber Technology Conference, London (1954) 397. Kutty S K N, Chaki T K and Nando G B, Polym Degrad Stab, 38 (1992) 187.

5 Composites of polychloroprene rubber with short fibres of poly(ethy1ene terephthalate) and nylon M ASHIDA

5. I

Introduction

The reinforcement of rubber compounds with short fibres has become necessary in many products, especially in the tyre, hose and belt industries. An advantage of short fibre-rubber composites is that they preserve the characters of both the elastic behaviour of rubber and the strength and stiffness of reinforcing fibres. 1 ~ Moreover, short fibres in the composites provide dimensional stability during fabrication and in extreme service environments by restricting matrix distortion, and improving creep re ~ist a nc eThey .~ also improve tear and impact strength' by controlling growth of internal crack^.^-^ The mechanical properties of the composites, such as modulus, strength and ultimate elongation, aepend on fibre orientation, fibre aspect ratio and adhesion between fibre and rubber matrix.' Moghe' reported on the milling parameters which cause fibre orientation and so influence the composite properties. According to Coran ef al," the properties of cellulose fibre-elastomer composites depend on the type of elastomer used as the matrix, the fibre concentration, fibre aspect ratio and fibre orientation. O'Connor ' compared composites reinforced with five kinds of fibres and found that their mechanical properties depend on fibre type, fibre content, fibre aspect ratio, fibre orientation, fibre dispersion and fibre-matrix adhesion. Blackley and Pike" reported the viscoelastic properties of the composites in a series of articles. Goettler and Shen13 presented a general review on short fibre reinforced elastomers. Abrate14 reviewed the mechanics of short fibre reinforced composites. This chapter presents the results of a study on the mechanical and viscoelastic properties of short fibre-chloroprene rubber composites and on the stress decay and stress relaxation of the composites under dynamic fatigue.

5.2

Preparation of composites

The elastomer used is chloroprene rubber (CR) and the recipe of rubber compounds is presented in Table 5.1. Short fibres used are Nylon 6 and

3

Composites of polychloroprene rubber

I17

Table 5. I . Compound recipe Ingredients

Parts by wt

CR Carbon black Process oil Stearic acid Antioxidant MgO ZnO Ethylene thiourea Nylon 6 fibre (diameter 27 pm) Length: P E T fibre (diameter 21 p n )

100 36 4 2 2 4

5 0.5 0.5-8.0 rnm

poly(ethy1ene terephthalate) (PET). Their lengths and diameters are presented in Table 5.1. In order to promote their adhesion to rubber matrix, nylon fibres were treated with a bonding agent ('RFL') and PET short fibres were dipped in isocyanate solution in advance and were then treated with RFL system. The RFL dip was formulated as follows:(quantities in wt%);' resorcinoll.89, formaldehyde(37%) 2.77, water 48.55, aqueous NaOH (1.67%) 3.06, styrene butadiene rubber (SBR)-vinylpyridine latex (41%) 41.78, concentrated NH,OH (28%) 1.95. The dry add-on of adhesive is approximately 9%. All the compounds and short fibres were mixed in a Banbury mixer and milled for a time sufficient to disperse the fibres in the matrix at a mill opening of 2 mm. The milling direction was always kept the same in order to give a maximum amount of short fibre alignment. Finally, each stock was passed through the mill to form a sheet of 2.0mm thickness. All samples were cured by compression moulding at 153"C for 20 min and were stored in dry air.

5.3

Mechanical and viscoelastic properties 5.3. I

Processing effects

Owing to the high viscosity of rubber compounds, short fibres are buckled and broken by high shear stress during the mixing process and their fibre length distribution is different from the original length.' L*'6-18 After the mixing of CR compounds with short fibres that are 6 mm in length, the distribution of fibre lengths is classified as one of three types depending on fibre species, as shown in Fig. 5.1. Short fibres of PET and vinylon retain their original length of 6mm without breakage (I). Nylon, aramid and rayon fibres are buckled or broken and give rise to a broad distribution in a range of shorter lengths (11). Carbon and glass fibres are broken into much smaller pieces and their lengths reduce to about 150 pm (111). By repeated passing through a mill, the short fibres in the composites tend to align along the flow direction and give an anisotropy to the composites.9~'sThe tensile stress required to produce respectively 10% and 20% strain in nylon 6

I18

Short fibrepolymer composites

50 40

30

20 10

0 0

2

4

length

6

8

(mm]

5.1 Length distribution of fibres after mixing

fibre-CR composites loaded with RFL treated and non-treated fibres which are 6mm in length is shown in Fig. 5.2 as a function of the angle 0 to the rolling direction. The relationship between the tensile stress and the rolling direction shows strong orientation of fibres along the flow direction within various composites." The results indicate both that short fibres assume high orientation in these composites and that the effect of RFL treatment appears in the composites loaded with over 5 vol% fibre. The swelling geometry of unidirectional composites is explained by Coran et uf" as follows. The length of a specimen I,, measured at an angle 0 to the fibre direction, expands to a length I , during swelling. The length (yo) parallel to the fibre direction changes to y, and the length (x,) transverse to the fibre direction changes to x, after swelling. If the swelling ratios for the transverse and longitudinal directions are uT = x,/xo and aL= y J y , respectively, then the swelling ratio at an angle 8 ( = a, = IJI,) is given by: u i = (u:

-

uE)sin28 + u t

L-5.11

The swelling of short fibre-CR composites in benzene also displays the anisotropy of the composites shown in Fig. 5.3.16 In the swelling the strong constraint parallel to the fibre direction shows the high orientation of fibres and

Composites of polychloroprene rubber

I I9

0

8 (degrees) 5.2 Tensile stress as a function of angle8 to the fibre orientation required to give 10% strain ( 0 , 0) and 20% strain (m, 0)in nylon 6 fibre-CR composites loaded with 6 m m long 1 0 ~ 0 1 %fibres and are RFL treated ( 0 ,m) and non-treated (0, 0).

the strong adhesion between the rubber and the, fibres. The straight lines of RFL-treated fibres, uz - u;, have the same slope for both composites, and the interceptsu; are 1.0 and 1.03for PET fibre-CR and nylon 6 fibre-CR composites, respectively. This discrepancy may be due to buckling of nylon 6 fibres during the mixing.

5.3.2 Mechanical properties Figure 5.4 shows the stress-strain curves of PET fibre-CR and nylon 6 fibre-CR composites, loaded with 10 vol% fibres that are 6 mm in length, in the longitudinal direction in comparison to the cured CR compound and to PET and nylon 6 filaments. The stress of composites loaded with a particulate filler increases along a similar curve to the cured CR compound. The stress in composites loaded with short fibres increases almost linearly with increasing strain, which is similar to the stress-strain behaviour of the fibres. The composites break at an elongation under 20% and 30% when loaded with PET fibre and nylon 6 fibre, respectively. The tensile strength of PET fibre-CR composites loaded with fibres less than 1 mm in length decreases almost linearly with increasing fibre loading, as shown

120

Short fibre-polymer composites

2.0

1.8

1.6

1.4

1.2

1.o 0

0.2

0.4

0.6

0.8

1.o

5.3 Relationship between %2 and sinZO for PET fibre-CR composites with RFL-treated fibres (a)and non-treated ones (0) and for nylon 6 fibre-CR composites with RFL-treated fibres (H)and non-treated ones (0).

in Fig. 5.5,20while on loading with fibres longer than 2 mm in length the tensile strength falls steeply to the minimum value at 5 ~ 0 1 %loading, then the plot reverses and the strength increases with further increase in fibre loading. The elongation at break falls gradually as the loading with short fibres that are 1 mm or less in length increases, which is a similar trend to the tensile strength. The elongation at break of composites loaded with fibres that are 2 mm or more in length decreases in the same way as on loading with the shorter fibres up to 5 ~ 0 1 %loading, but at loadings over 5 ~ 0 1 %i t falls markedly and an inflexion point appears corresponding to a change at the minimum value of the tensile strength. The effect of fibre length on the stress-strain curves for PET fibre-CR composites loaded with 10 vol% fibre is shown in Fig. 5.6. For composites loaded with PET fibres from 4 to 8 mm in length, the tensile stress increases monotonically with strain until they break at an elongation under 20% at about 20 M P a stress. For composites loaded with PET fibres that are 1 mm or less in length, the tensile stress increases gradually with increasing strain until the yield point at an elongation of about 20%, and then it approaches a plateau. In contrast, the composite loaded with 2mm PET fibres breaks soon after the yield point.

I21

Composites of polychloroprene rubber

1.2

30

1.o

c5

m

p

0.a

20

Y

0.6 v) v)

Q)

L

0.4

10

P

CI v)

0.2

I

0

0 0

10

20

strain

30

1%)

Stress-strain curves for cured CR (-), PET filament (-----), nylon 6 filament (-----), PET fibre-CR (-.-.-) and nylon 6 fibre-CR (-..-..) com-

5.4

posites. Therefore, it seems that 2 mm is the critical fibre length to reinforce this class of composites. The stress-strain curves for PET fibre-CR composites loaded with different concentrations of 2mm fibres are shown in Fig. 5.7. The stress of the composite loaded with 5 vol% fibre increases gradually as the strain is increased and, after yielding at an elongation of about 28%, a nearly constant stress is maintained similar to the behaviour of the compositescontaining 10 vol% offibres 1 mm long orshorter.On theotherhand,forthecompositesloadedwith fibretoover 1 0 ~ 0 1 % the tensile stress increases monotonically with increasing strain until failure occurs. As the loading is increased, the tensile stress increases more steeply and the composites break at a smaller elongation and at a higher stress. Furthermore, it is difficult to mix short fibres longer than 4 mm in length with rubber in the mill and the fibres are dispersed very poorly in the composite, but short fibres of 2 mm in length are easily mixed up to 25 vol% and are dispersed well in the composite. The effect of fibre length on the tensile stress required to produce a given extension is shown in Fig. 5.8 for composites loaded with 1 0 ~ 0 1 %PET fibre. The tensile stress of composites in the direction perpendicular to the fibre direction at an elongation of 100% (Mloo) increases considerably as the fibre length is increased up to 2 mm, after which it remains nearly constant as the fibre length is increased from 2 to 8 mm. The tensile stress of composites in the direction of fibre orientation at an elongation of 10% (Mlo)increases rapidly and almost linearly as the fibre length is increased up to 2 mm, but with fibres of 2 mm or more in length the tensile stress increases much more slowly with fibre length. In the

I22

Short fibrepolymer composites

25

I---

'

20

-

\ .-

'

/

//!

CI

P,

c

Q) L

15

10

0

1

I

I

I

I

2

4

6

8

10

5.5 Effect of fibre loading (Vf) on tensile strength of PET fibre-CR composites loaded with fibres of 0.5 mm (O), 1 rnm ( O ) ,2 rnm ( A ) , and 4 rnrn (V)in length.

direction perpendicular to the fibre direction, short fibres in the composites act in a manner similar to that of fillers such as calcium carbonate or clay; that is, the fibres have a reinforcing effect on the matrix rubber by acting as massive crosslinks.21 . 2 2 It is noted that the mechanical properties of PET fibre-CR composites change non-linearly at a fibre length of about 2 mm. Consequently, the CR composites reinforced with short fibres give rubber-like or fibre-like behaviour, depending on the fibre loading and the fibre length. 5.3.3 Dynamic properties23 The effect of the volume fraction of fibres on the storage modulus E' and loss modulus E" in the fibre direction is shown in Fig. 5.9 for PET fibre-CR composites loaded with fibres of 6 mm in length. The storage moduli of PET fibre and nylon 6 fibre are higher than that of the cured CR compound by a factor of more than lo2 over the whole temperaturerange above 0 "C. The rubbery state of cured CR compound is in the range 0-160°C. The composites have a higher storage modulus that is about 20 times higher than that of the cured CR compound and exhibit the rubbery state in the range 0-lOO"C, in which the storage moduli of the composites decrease linearly with rising temperature. The

I23

Composites of polychloroprene rubber

25 20 n

Y

15 10

5

0 0

10

20

30

40

strain (%) 5.6

Stress-strain curves of PET fibre-CR composites loaded with 10 vol% fibres of various lengths.

storage modulus of the nylon 6 fibre-CR composites displays similar behaviour, but the modulus of the composite loaded with 15vol% fibre has the same value as the composite loaded with 12.5 vol% fibre, because nylon 6 fibres tend to buckle during mixing and so a uniform dispersion of the fibres is not obtained in the matrix. Therefore, the effect of increasing fibre loading for nylon 6 fibre-CR composites does not appear when the loading is over 12.5 ~ 0 1 % . The loss modulus of cured CR compound shows a sharp peak at - 33 "C, whereas two peaks of the loss modulus appear at - 20,120 "C,and - 20,90 "C for PET fibre and nylon 6 fibre, respectively. Both values of E' and E" increase to higher values with increasing fibre loading. In addition, the storage modulus of PET fibre-CR composites given in Fig. 5.10 as a function of fibre length shows that in the rubbery region the storage modulus of composites loaded with fibres shorter than 1 mm in length has a similar concave curve to the CR compound, and that the storage modulus of composites loaded with fibres longer than 2 mm in length gives straight lines. For the composites loaded with fibres that are 2 mm or less in length, the elastic modulus increases linearly as the fibre loading is increased, as shown in Fig. 5.10, and the slope increases with the fibre length up to 2mm and is constant over the range of fibre length from 4 to 8 mm. The relationship between storage modulus and fibre orientation is represented as follow^:^

I 24

Short fibre-polymer composites

30 I

I

I

I

I

1

25%

c

a

2

20

Y

I / / /Y 10%

10

0 0

10

20

strain

30

40

(XI

5.7 Effect of fibre loading on stress-strain curves of PET fibre-CR composites loaded with 2 mm long fibres. Dashed line is the composite loaded with 6 mm long 10 vol% fibres. (Loadings given in vol%.)

where EL is the modulus of composite wherein the fibre direction deviates from the extensional direction by the angle 9, EL is the modulus along the direction parallel to the fibre alignment (9 = 00) and E; is the modulus in the transverse direction to the fibre direction (0 = 90"). The E' of the composite decreases as the angle ( 0 ) increases and the lowest value is in the transverse direction (0 = 9 0 O ) . Figure 5.1 1 shows that the measured moduli agree well with the calculated values for nylon 6 fibre-CR and PET fibre-CR composites. The temperature dependence of loss tangent tan6 for nylon 6 fibre-CR and PET fibre-CR composites loaded with 6 mm long 10 vol% fibres and cured CR compound is shown in Fig. 5.12. A large dispersion peak is observed at -28 "C for the cured CR compound. In the case of dynamic modulus of fibres, two moderate peaks are observed at - 16, 100 "C and - 16, 135 "C for nylon 6 and PET fibres, respectively; the peak at the lower temperature corresponds to the P-dispersion, whereas the peak at the higher temperature corresponds to the a-dispersion ascribed to micro-Brownian motion of amorphous chains.24 For nylon 6 fibre-CR and PET fibre-CR composites, each p-dispersion is observed as a very small shoulder at the corresponding temperature - 16 "C, because the p-dispersion peaks of fibres are weak compared with that of CR. Therefore, the peak at - 28 "C corresponds to the main dispersion of C R and the peak at the

I25

Composites of polychloroprene rubber

15

10

5

i

0' 0

I

I

I

1

2

4

6

8

fibre length

(mm)

5.8 Effect of fibre length on tensile moduli for PET fibre-CR composites loaded with 10~01% fibres: stress measured parallel to the fibre direction at 10% extension ( 0 ) ;stress measured normal to the fibre direction at 100% extension (0).

higher temperature corresponds to the u-dispersion of each fibre. In addition to these two peaks, a small and broad peak is observed at about 90 "C for the PET fibre-CR composite. The additional tan ?I peak decreases with increasing fibre length. The temperature dependence of tan 6 for short fibre-CR composites displays more clearly the effect of the fibre orientation. As the angle to the fibre direction increases, the peak at - 28 "C rises progressively and another peak at the higher temperature falls to lower values. For PET fibre-CR composites the broad peak at 90 "C protrudes with increasing the angle and becomes a relatively sharp peak in the transverse direction to the fibre orientation. Such a dispersion peak is not observed for the nylon 6 fibre-CR composite, probably because the peak is shielded by the a-dispersion of nylon being in this temperature range. The relationship between the storage modulus at 20 "C and the volume fraction of fibres for PET fibre-CR composites shows that the storage moduli increase linearly with increasing fibre loading and take same value for fibre lengths from 4 to 8mm, as shown in Fig. 5.13. It is well known that the mechanics of reinforcement of short fibre-elastomer composites has been predicted by the Halpin-Tsai Unfortunately, however, the moduli given by the

I26

Short fibre-polymer composites

1o4

-10’

a

e

I

Y

1o2

\\

W . )

\

wlo t

1 -100

I

I

0

100

temperature

200

(‘C)

Effect of fibre loading on E‘ and E” for PET fibre-CR composites loaded with 6mm long fibre: fibre contents are 5 ~ 0 1 % (V,Y),7.5~01%( 0 ,+), 10~01%( A , A),12.5~01%(0, m) and 15~01%(0, 0). 5.9

Halpin-Tsai equation are lower than those obtained experimentally.

5.3.4 Reinforcing mechanism The storage modulus of composites is given by the parallel model as follows:

E:

== E;Vf

+ ELV,

= (E;-

E;)Vf

+ EL

C5.31

where E:, E; and E L are the storage modulus values for composite, fibre and matrix and V, and V, are the volume fractions of fibre and matrix, respectively. This ‘parallel model’ is based on the assumption that matrix and fibres are strained to the same extent.” The findings suggest that the fibres are bonded strongly to the matrix so that the fibre strain equals the matrix strain in the range of the tensile deformation which is applied to the composite by the viscoelastometer. Consequently, if equation [5.3] is modified as follows

E,

= x(E; - E’,)Vf

+ EL

C5.41

where M. is a coefficient depending on fibre length, equation [5.4] can be applied well to these short fibre-CR composites. The relationship between the length of short fibres and the storage modulus of

I27

Composites of polychloroprene rubber

10'

1

I

1o3 n

(II

e

2 -1

o2

\

10

L-

'

CR matrix 1 '

-100

I

I

0

100

temperature

200

('C)

5.10 Effect of fibre length on storage moduli for PET'fibre-CR composites loaded with 10 vol% fibres: fibre lengths are 0.5 mm (o),1 mm (A), 2 mm (0) and 4 8 mm (0).

-

the composite at 20°C is shown in Fig. 5.14. For a fixed loading, the storage modulus increases linearly as the fibre length is increased and at a length of 4 mm or above the elastic modulus attains a constant value which increases as the loading is increased. The coefficient a in equation [5.4] is obtained from the observations.'' Consequently, the storage modulus of the short fibre-CR composites is given by equation C5.31 when the fibre length (1) is longer than 3 mm and when I is shorter than 3mm it is given by:

Ek

=

EL

+ 0.341(E; - EL)V,

C5.51

The storage modulus of a short fibre-CR composite can be obtained from the storage moduli of the component materials, the fibre length and the fibre loading by using equation [5.3] or C5.51. However, the tensile stress differs according to the diameter of the fibre^'^,'^ and so it suggests that the coefficient a depends on the elastic modulus and the diameter of the short fibres used. It is noted that the mechanical and dynamic properties of short fibre-rubber composites change non-linearly at a fibre length around 3 mm.

Short fibrepolymer composites

I28

400

-

300

m

200 \

W

100

v

-

0

30

60

90

(degrees)

8

5.11 Effect of fibre orientation on storage moduli for nylon 6 fibre-CR (0) and PET fibre-CR composites ( 0 )loaded with lOvol% fibres. The storage

moduli calculated from equation [ 5 . 2 ] are represented by the lines.

On the other hand, the storage moduli (E:,T)of short fibre-CR composites at 20 “C measured in the direction transverse to the fibre orientation are much larger than those calculated from the Halpin-Tsai e q ~ a t i o n . ~ ’ ~ Log ’~ increases linearly with increasing fibre loading for PET fibre-CR composites, as shown in Fig. 5.1 5.27 The value of log Eb,Tis expressed as the following equation: log Eb,T

= M.

+ p Vf

whereor and b are constants. The values ofor and b are obtained from the intercept and the slope of the experimentally obtained lines, respectively. It is well known that the moduli of many composites composed of two phases can be expressed by ‘the logarithmic law of mixing’ as the following equation:29

log EL

=

V,,, log E:,

+ V, log E; = log E L + Vf log(E;/Eb)

~5.61 The values a agree with logEk for all composites and the values of p of the composites with untreated PET fibres are nearly equal to the values of log E;/E;. Therefore, equation C5.61 can be applied well to the transverse storage modulus of those short fibre-CR composites loaded with untreated fibre. In the case of RFL-treated PET fibres, the slopes are larger than log E;/E:, by a factory giving: log E:.T = log E:,

+ y v,log(E;/E:,)

C5.71

I29

-100

0

temperature 5.12

100

200

('C)

Temperaturedependence of tan 6 for cured CR compound(O ) ,nylon 6 fibre-CR (0)and PET fibre-CR composites (A).

where y is a factor depending on the character of the interphase between fibres and matrix, that is, the degree of the bonding force. If the fibre is not treated with RFL solution, they value is unity because no interphase is formed between fibres and rubber matrix, and log EL,Tincreases with increasing fibre loading according to the volume effect. When the fibre is treated with RFL solution, the apparent volume fraction of fibre becomes larger than the true volume fraction of fibre with increasing that is, y Vf is larger than Vf and y is larger than unity. The modulus of particle filled elastomer composites has been given by 'the logarithmic law of mixing'31 or the modified Gray and M ~ C r u m have ~~ indicated that the modulus of a semi-crystalline polymer treated as a two-phase composite consisting of the crystalline region and the amorphous one is also given by equation C5.61. Therefore, it seems that short fibres in the transverse direction act just like particles or crystalline regions to reinforce the matrix rubber. The peak values of tan 6,at - 28 "C correspond to the main dispersion of the

Short fibrepolymer composites

I30

? 50

100

ti

H

1

\

w 50

0 0

2

4

6

8

10

5.13 Effect of fibre loading on storage modulus for PET fibre-CR composites loaded with fibres that are 0.5mm fo),t m m (01, 2 m m (A)and 4 8 mm (-----) in length.

-

cured CR compound. Figure 5.16 shows the relationship between the peak values of tan6 at -28 "C and the volume fraction of fibre for the PET fibre-CR and nylon 6 fibre-CR composites loaded with fibres that are 6 mm in length. Similar to other composites,28 the peak values of tan 6c,T for samples with their major axis transverse to the fibre direction decrease with increasing fibre volume fraction and are represented by the following equation: tan 6c,T= tan 6,, - a V,

~ 5 . ~ 1

where tan6,,, and tan6, are the loss tangent for composite and matrix, respectively, and a is a coefficient depending on the fibre type. If the mechanical damping of composites comes only from the matrix polymer, the damping ratio of the composite to the matrix is roughly equal to the volume fraction of p ~ l y m e r .Namely, ~ ~ ? ~ ~the peak values of tan should decrease in proportion to volume fraction of fibre along the dashed line shown in Fig. 5.16. In the case of the composite loaded with untreated fibres, the peak values of tan hc,Tare on an identical line, whose slope is steeper than that of the dashed line. Furthermore, the peak values of tan 6c,Tfor the composites with RFL-treated fibres decrease along the steeper lines, which depend on the fibre type, with increasing fibre loading. The results indicate that a region with strong interaction is formed between the

131

Composites of polychloroprene rubber

2

0

4

fibre length

6

8

(mml

Effect of fibre length on storage modulus at 20°C for PET fibre-CR composites loaded with fibres of 1 ~ 0 1 %(O), 2 . 5 ~ 0 1 %(O), 5 vol% ( A ) and 5.14

lOvo1?4, (V).

fibres and the CR matrix and gives rise to a decrease in the apparent volume fraction of the matrix. The dispersion peak of tan6 appearing at 90°C for the PET-CR composite may be caused by the interface region in the composite. On the other hand, the peak value of tan 6 at - 28 "C for longitudinal samples is considerably lower than that for transverse samples and decreases parallel to the dashed line with increasing fibre loading. As mentioned previously, since the dynamic modulus in the longitudinal direction is strongly affected by the fibres, the peak values of tan 6 at - 28 "C probably depend not only on the volume fraction of the matrix but also on the properties of the fibre.

5.4

Effect of absorbed Waterj6

Water absorbed by hydrophilic fibres such as nylon and rayon has a profound effect on the mechanical properties of these fibres. The effect of absorbed water on the dynamic mechanical properties of various polyamides has been s t ~ d i e d . ~ ' * ~ * Physical properties of polyesters are considerably affected by absorption of a small amount of ~ a t e r . ~ Willett41 ~ . ~ ' has reported the effect of temperature, humidity and geometric structure on the viscoelastic properties of tyre cords such as rayon, nylon-6,6 and PET. Figures 5.17 and 5.18 show the temperature

Short fibre-polymer composites

100

10

0

5

10

15

v, (%I ElTect of fibre loading on transverse storage modulus, E;, at 20 "C for PET fibre-CR composites. The dotted line is calculated from equation C5.71,

5.15

wheny

=

1.

dependence of the dynamic storage modulus and loss modulus in the longitudinal fibre direction for nylon 6 fibre-CR composites, loaded with 6 mm long 10 vol% fibre, which contain water absorbed in a moist atmosphere. The storage modulus of the composites decreases in two steps with rising temperature. The first drop, appearing at about -40 "C,coincides with the glass transition of the CR matrix and the p-dispersion of nylon 6 fibre and occurs in the same temperature range independent of the absorbed water content. The second drop, appearing at about 60 "C, shifts to a lower temperature with increasing absorbed water content, and this step finally disappears when the composite absorbs 4.26% water. The loss modulus of the composites has two maximum dispersions: one dispersion peak at - 35 "C is caused by theglass transition of the CR matrix and the p-dispersion of nylon fibre and the other dispersion at 82 "C corresponds to the a-dispersion of nylon fibre. The strong low temperature peak does not change its height and position when the water content changes, but the medium strength higher temperature peak shifts to lower temperatures while maintaining its height when the absorbed water content increases. The tan 6 peak shifts to lower temperatures in a similar way to the loss modulus with increasing absorbed water. Becker and O b e r ~ have t ~ ~reported that the maximum of tan 6 for polyamide with 8% water content shifts to about 0°C. According to

I33

Composites of polychloroprene rubber

1.0 I

I

I

I

I

I

0

5

10

I

0.8 Y

-a id

c

0.6

4

0.4

0.2

15

5.16 Effect of fibre loading on the maximum tan6 at -28 "C,(tan 6),,,, for non-treated fibre-CR (0 = 90" (A)) and RFL treated PET fibre-CR (0 = 0" 8 = 90" ( 0 ) and ) nylon 6 fibre-CR (0 = 0 ( O ) ,0 = 90 (m)) composites.

(o),

O

Woodward et ~ 1 the, a-dispersion, ~ ~ appearing at about 100 "C for nylon fibre, shifts to lower temperatures with increasing water content and the peak reaches 0 "C for the 6.4% water specimen. The change has been attributed to segmental motion in the amorphous regions as a consequence of the breaking up of hydrogen bonds between chains4' Therefore, the a-dispersion of nylon fibre in the composite shows the same behaviour as that of fibre alone with increasing absorbed water. A small hump appears at 90°C after displacement of the a-dispersion peak to lower temperatures. The small hump in the dry composite is shielded by overlapping the a-dispersion peak of nylon fibre. The additional dispersion may be caused by a third component in the composite because the hump appears only in the composite. The temperature dependence of tan6 for the PET fibre-CR composite also shows little effect of absorbed water. The maximum peak at 135 "C corresponding to the a-dispersion of PET fibre shifts slightly to lower temperatures with increasing absorbed water content. A small shoulder appearing at about 90 "C diminishes as absorbed water increases and disappears for the composite containing 3.4% water. This dispersion corresponds to the additional dispersion of nylon 6 fibre-CR composite mentioned above. The tan 6 peak for the nylon 6

I34

Short fibre-polymer composites

n

103

\

W

1o2

10 -100

0

100

temperature

200

('C)

5.17 Effect of absorbed water on storage modulus as a function of temperature for nylon 6 fibre-CR composites: water contents are 0% (O), 0.13% (A), 0.42% (0)and 4.26% (0).

fibre-CR composite shifts at the approximate rate of 25 "C/o/, water. The rate is faster than those for nylon reported by Howard and Williams4* and Willett.4' On the other hand, the temperature shift of the PET composite is too small to determine the rate. These results indicate that nylon 6 fibre absorbs much more water than the CR matrix, while the water absorption for PET fibre is considerably less than for nylon 6 fibre. When the wet sample is heated for 1 h at 100 "C, the a-dispersion peak remains at 45 "C in spite of the fact that the absorbed water content decreases to 0.7%. After treating the wet sample for 1 h at 120 "C, the a-dispersion peak returns to 90 "C, which is the peak temperature of the original dry sample. The results suggest that nylon fibre absorbs a larger amount of water than CR matrix, while the water absorption of PET fibre is considerably less than that of nylon fibre. The absorbed water in nylon fibre is bonded more strongly than that in CR matrix, and its concentration is only slightly diminished by heat treatment under 100 "C, whereas the absorbed water in CR matrix is completely removed by the heat treatment.

I35

Composites of polychloroprene rubber

103

1

1 -100

1

I

I

I

I

0

100

temperature

I 200

('C)

5.18 Effect of absorbed water on dynamic loss modulus as a function of temperature for nylon 6 fibre-CR composites: water contents are 0% (0), 0.13% (o), 0.42% (A)and 4.26% (0).

5.5 Dynamic fatigue of composites Short fibre reinforced rubber has been applied to industrial rubber products such as conveyor belts, power transmission belts and hoses because of its advantageous properties. The power transmission belt represents a typical dynamic application for short fibre reinforced rubber. When a power transmission belt runs over pulleys, short fibre reinforced rubber in the belt receives complex deformation due to repeated tension, compression and bending at high speed. When short fibre-rubber composites are subjected to a dynamic fatigue test under cyclic strain or cyclic load, the changes in stress and strain are influenced by the short fibres, the rubber matrix, and the interfacial region between fibres and the rubber matrix.

5.5. I

Extensional fatigue

The specimens for the extensional fatigue test were cut out in a long plate of 100 x 20 x 5mm. The specimens of PET fibre-CR composites, loaded with 1 0 ~ 0 1 %of 6mm fibres, were extended parallel to the fibre orientation under repeated constant displacement to 5% of static strain and & 2.5% of sinusoidal strain.43The stress of composites during fatigue was obtained at the maximum strain of 7.5% from the Lissajous diagram. Figure 5.19 shows the stress and

I36

-

Short fibre-polymer composites

15

-

0 - 6 040

0

I

t

20

1o2

1o3

1o4

cycles 5.19

Stress decay and surface temperature of PET fibre-CR compositcs under repeated strain cycles.

surface temperature of composites during the fatigue tests a t 5 and 25Hz frequency. At a frequency of 25 Hz, the stress decreases rapidly from 100 cycles t o SO00 cycles and after that the composite registers a constant stress until 30 000 cycles. T h e surface temperature of the composite rises steeply from 250 cycles to 5000cyclcs and reaches the highest value at 7500 cycles, then the temperature falls slowly until 10 000 cycles, which corresponds to the stress decay of the composite under fatigue. At a frequency of 5Hz, the stress decreases gradually and the surface temperature rises a little. In the case of the cured C R compound, the surface temperature rises slightly after 1500 cycles. Because the rubber matrix can be regarded as a macroscopically homogeneous material which generates little heat under fatigue, the rising temperature of the composite is caused by differential displacements among the three components of fibre, interfacial region and rubber matrix. Under the high frequency, the interfacial region between the fibres and the rubber matrix is destroyed in 5000 cycles and much heat generates in the composite due to rapid displacement. The stress decay from 100 cycles to 5000 cycles becomes steeper, with increasing fibre loading for PET fibre-CR composites. In the case of nylon 6 fibreeCR composites, the rising temperature of the composite loaded with 15 vol% fibre is the same as the composite loaded with

Composites of polychloroprene rubber

I37

1 0 ~ 0 1 %fibre, because there is too much fibre in the former composite for good dispersion. It seems that some parts of the interfacial region of the composite loaded with 15 vol% fibre are not subjected to sufficient deformation to affect heat generation of the composite. The surface temperature of the composite loaded with 5 vol% fibre rises in the same manner as with composites loaded with 10 and 15 vol% fibres, although the former gives temperatures lower than the latter over the whole range. This is because the composite loaded with 5 vol% fibre generates less heat, owing to the small interfacial region. The stress decay probably affects the destruction of the interfacial region and the internal heat generation of the composites. At about 5000 cycles, the interface region is broken away and short fibres start to slip in the composite, and the temperature of the composite falls gradually. The destruction of the interface region is completed after 5000 cycles and the stress remains at a constant value, because the strain amplitude is always constant. Consequently, the stress decreases significantly and the surface temperature rises before the destruction of the interfacial region. After the destruction no heat is generated in the composite since short fibres slip in it. The stress decreases significantly and the surface temperature increases with increasing dynamic strain. The stress decay in dynamic fatigue may be affected by the destruction of the interfacial region and the rising temperature of the composite. When the dynamic fatigue test is carried out at temperatures higher than the maximum temperature reached in the fatigue test in ambient air at 25 "C, the stress decay of the composites shows the same tendency as that of the composites at 25 "C, although the initial stress decreases with increase in t e m p e r a t ~ r eIt. ~seems ~ that the stress decay of the composites in dynamic fatigue is caused mainly by the destruction of the interfacial region between the fibres and the CR matrix. Figure 5.20 shows the results of fatigue tests carried out discontinuously, that is, three runs of a pause time of 400 s after 10 000 cycles in the same extensional condition as shown in Fig. 5.19. The stress after the first pause of 400 s recovers to a value higher than before the pause and at the same time, the rising temperature in the first run falls to the initial temperature. The stress after the first pause of 200 s is similar in value to that after a pause of 400 s. In the second and third runs the stress decreases linearly with increasing number of cycles after both pauses of 200 s and 400 s. When the pause time is as short as 10s, the surface retains the high temperature rise from the first run, and the difference in stress before and after the pause is very small. It seems that the stress decay of the composites is caused not only by destruction of the interfacial region between the fibres and the rubber matrix but also by the rising temperature of the composite. The results suggest that the composites are able to recover the strain including the interfacial region between fibres and the rubber matrix, if the composite should be given enough pause time before the destruction of the interfacial region in repeated cycles of strain. The surface temperature of the cured CR compound shows isothermal contour lines which are symmetrical about both horizontal and perpendicular axes of the specimen.45 On the other hand, the surface temperature of short fibre-CR

Short fibre-polymer composites

I38

15 0

10

5 60

t

L

I-

1 s t -I

I---

2nd -I

-I

3rd-I

25Hf

cycles 5.20 Effect of frequency on stress decay and surface temperature of PET fibre-CR composites in discontinuous fatigue tests.

composites loaded with 10 vol% fibre increases steeply from the edge towards the centre part of the specimen under dynamic fatigue. These results indicate that the unfilled rubber deforms homogeneously, while short fibre-CR composites deform heterogeneously owing to the existence of short fibres. F o r the composite loaded with 10 vol% fibre, the maximum temperature after 10 000 cycles is higher than after 6000 cycles and the temperature contours after 10000 cycles become more closely spaced at the centre part and more broadly a t the extremity than those after 6000 cycles. The P E T fibre-CR composites loaded with 6 m m long 10 vol% fibre have been extended in various directions from parallel (angle = 0") to perpendicular (angle = 90") to the fibre direction.44 When the specimen is extended at 15" to the fibre direction, the stress decreases steeply until 5000 cycles and maintains the same value from 5000 cycles to 30000 cycles, which is similar to the case of longitudinal extension. O n extension at 30" to the orientation axis, the stress changes slightly until 5000 cycles. O n the other hand in the case of the directions of 45" and 90" to the fibre direction, the stress decreases slightly, apparently linearly. Figure 5.21 shows that the rise in temperature becomes lower with increasing angle to the fibre direction.44 The symmetry axis of the temperature distribution makes higher inclination to the extensional direction as the exten-

Composites of polychloroprene rubber

I39

150

15' 30'

45'

90'

0 1 1o2

I

I

10'

I 04

1o5

cycles 5.21 Surface tempcratures of cured CR compound and PET fibre-CR composites loaded with 10 vol% fibre during repeated extensional fatigue The extensional direction of specimens is at angle 0 to the fibre direction.

sion angle increases. At the angle of 90", the temperature distribution in the composite is fairly similar to that in the cured CR compound. The stress and the surface temperature of the composite gradually approach those of the cured CR compound as the angle to the fibre direction increases. These results suggest that the stress decay and the rising temperature are caused by shear deformation of the interfacial region along the fibre direction. The strain of the composites changes according to one of two cases under a repeatingconstant load such as 6 f 2,8 k 2,8 f'4 and 10 -t 2 M Pa ofsinusoidal In one case the strain increases a little and linearly with an increasing number of cycles and the initial strain and the increment of strain are in inverse proportion to the fibre loading. Both the initial strain and the strain increment of composites loaded with nylon 6 fibre are larger than the composite loaded with PET fibre. In the other case, appearing at higher loads and lower fibre contents, the strain increases remarkably with an increasing number of cycles and the strain of the composite loaded with PET fibres is larger than the composite loaded with nylon 6 fibres. Therefore, it seems that the former case is caused by deformation of the interfacial region between the fibre and the rubber matrix and

I40

- -

Short fibrepolymer composites I

10%

n

10

II

2

U

v) v)

a L

5

c. v)

0

0

strain

60

time

120

180

(s)

5.22 Stress-strain and stress relaxation curves of PET fibre-CR composite loaded with 1 0 ~ 0 1 %fibre:strain ratesare4%/min(-), S%/min(-.---.)and 400%/min (-----).

in the latter case deformation takes place by the destruction of the interfacial region. The change in strain in the latter case is restrained by the R F L treatment of fibres and by the high content of short fibres. These results suggest that the interfacial region of nylon 6 fibre-CR composite is so tenacious that it deforms to a large scale. In the case of the same load under the dynamic fatigue of larger displacement, the destruction of the interfacial region occurs in the early stages of the dynamic fatigue, because the interfacial region hardly deforms with sinusoidal stress. The stress-strain and stress relaxation curves of PET fibre-CR composites loaded with 1 0 ~ 0 1 %fibre are shown in Fig. 5.22 under 10 MPa of initial stress and 4, 80 and 400%/min of initial strain rates4' The stress-strain curves show that the faster the initial strain rate, the higher is the modulus of composites. This phenomenon seems to be caused by delay of flow of the interfacial region between fibres and the rubber matrix under the deformation process. On the other hand, the faster the strain rate, the larger the stress decay at the initial stage. It suggests that the interfacial region suffers large distortion under a fast strain rate. The stress relaxation of the PET fibre-CR composite after swelling in toluene suggests that the composite includes a weak bond between fibres and the rubber matrix, because the stress relaxation of composites after swelling is smaller than before swelling. From observations of the fracture surfaces of composites,

-I

-

141 1

I

I

5%

-

-----____

-

I

0

strain

60

time

120

180

($1

Compression stress-strain curves and compression stress relaxation curves of PET fibre-CR (-) and nylon 6 fibre-CR(----) composite loaded with 5 vol% and 15 vol% fibres: initial stress, 2 MPa.

fracture is caused by propagation from a nucleus, which is initiated by the destruction of the interfacial region surrounding a short fibre.46

5.5.2 Compressive fatigue4' The specimens for the compressive fatigue test were cut into cubes of 25 mm in length and were compressed in either the fibre direction or transverse to it by an applied stress of 2 or 4 MPa and at compression rates of 1.5 and 15 mm/min. Short fibre-CR composites have a higher modulus than the rubber matrix depending on the fibre loading and the compression direction. The higher the fibre loading the higher the modulus, and composites compressed in the fibre direction show a higher modulus than those compressed in the transverse direction. In the case of compression in the fibre direction, the high modulus of composites may be caused by forces in the short fibres and in the interfacial region between fibres and the rubber matrix against the compression. In the case of compression transverse to the fibre direction, the composites are easily deformed by a light compression like the cured CR compound. PET fibre-CR composites have a higher modulus than nylon 6 fibre-CR composites. This means that the stiffness and orientation of PET fibres are higher than those of nylon fibres. The compression stress-strain curves at a compression rate of 1.5 mm/min and the compression stress relaxation at the initial stress of 2 MPa are shown in Fig.

I42

Short fibre-polymer composites

5.23 for nylon 6 fibre-CR and PET fibre-CR composites. The strain ofcomposite loaded with 5 vol% fibre is about twice as large as that of the composite loaded with 15 vol% fibre at the initial stress of 2 MPa. At the initial stage the stress decay of the composite loaded with 5 ~ 0 1 %fibre is more rapid than for the composite loaded with 15 vol% fibre. After this, the stress decay of the former shows the same tendency as the latter. This phenomenon indicates that the stress decay of the composite loaded with 5 vol% occurs remarkably at the initial stage because the flow and the destruction of the interfacial region are caused by the deformation under compression. The stress decay of nylon 6 fibre-CR composites is more rapid than that of PET composites. It seems that the interfacial region between nylon 6 fibres and CR matrix is more flexible and more tenacious than that between PET fibre and CR matrix, as mentioned previously in the extensional fatigue test. The stress decay of the composite compressed at 15 mm/min is faster than that compressed at 1.5 nim/min. The stress decay of the composite loaded with 5 vol% fibre is larger than that of one loaded with 15 vol% at 1.5 mm/min of compression rate. O n the contrary, at 15 mm/min compression rate, the stress decay of the composite loaded with 5 vol% fibre is smaller than that loaded with 15 vol% fibre. In the case of compression with high strain rate, the composite absorbs insufficient energy to recover the deformation and the interfacial region stores the distortion in the compression process. At the initial stage of the stress relaxation, the stress decay of the composite arises remarkably because the interfacial region releases the distortion rapidly. Consequently, the shear deformation of the interfacial regions affects the stress and the surface temperature of the composites under compressive fatigue. If the fibre in the composite has low tenacity, the interfacial region of the composite is subjected to low shear deformation. Acknowledgement

The author wishes to thank his collaborators, S Mashimo, Y Yamaguchi, M Nakajima and T Noguchi in Mitsubishi Belting Ltd, for their special assistance. He is also indebted to W G u o for his kind help in preparing the figures. References 1 2 3 4 5 6 7 8 9 10

Boustany K and Hamed P, Rubber World, 171 (1974) 39. Beatty J R and Hamed P, Elastornerics, 110 (1978) 27. Rogers J W, Rubber World, 183 (1981) 27. Turner S, Br Plast, 38 (1965) 44. Beatty J and Hamed P, Elastornerics, 110 (1978) 27. Lavengood R E and Gulbransen B, Polym Eng Sci, 9 (1969) 365. Hangerup E, J Appl Polym Sci, 7 (1963) 1093. Derringer G C, J Elastoplastics, 3 (1971) 230. Moghe S R, Rubber Chem Technol, 49 (1976) 1160. Coran A Y , Boustany K and Hamed P, Rubber Chem Technol, 47 (1974) 396.

Composites of polychloroprene rubber

I43

11 O'Connor J E, Rubber Chem Technol, 50 (1977) 945. 12 BlackleyDCand PikeNT, KnutschukGummiKun.~tstofli.,29(1976)607,680;30(1977) 367; 31 (1978) 16. 13 Goettler L A and Shen KS, Rubber Chem Tcchnol, 56 (1983) 619. 14 Abrate S, Rubber Chem Trcimol, 59 (1986) 384. 15 Foldi A, Rubber Chem Technol, 49 (1976) 379. 16 Noguchi T, Ashida M and Mashimo S, Nippon Gomu Kyokaishi, 56 (1983) 768. 17 Noguchi T, Ashida M and Mashimo S, Nippon Gomu Kyokaishi, 57 (1984) 171. 18 Noguchi T, Ashida M and Mashimo S, Nippon Gomu Kyokaishi, 57 (1984) 829; Int Polym Sci Technol, 12 (1985) T67. 19 Coran AY, Boustany K and Hamed P, J Appl Polym Sci, 15 (1971) 2471. 20 Ashida M, Noguchi T and Mashimo S, Nippon Gomu Kyokaishi, 60 (1987) 158;Compos Polym, 2 (1989) 339. 21 Ashida M, Noguchi T and Mashimo S, J Appl Polym Sci, 30 (1985) 1011. 22 Souma I, J Appl Polym Sci, 27 (1982) 1523. 23 Ashida M, Noguchi T and Mashimo S, J Appl Polym Sci, 29 (1984) 661. 24 Dumbleton J H and Murayarna T, Kolloid 2 2 Polym, 220 (1967) 41. 25 Halpin J C, J Compos Muter, 3 (1979) 732. 26 Halpin J C and Kardos J L, J A p p l Phys, 43 (1972) 2235. 27 Nielsen L E, J A p p l Phys, 41 (1970) 4626. 28 Lewis T B and Nielsen L E, J Apply Polyrn Sci, 14 (1970) 1449. 29 Nielsen L E, Mechanical Properties of Polymer and Composites, Marcel Dekker, New York (1975). 30 Ziegel K D, J Colloid Interf Sci, 29 (1969) 72. 31 Fujirnoto K and Nishi T, Nippon Gomu Kyokaishi, 43 (1970) 54. 32 Nohara S, Kobunshi Kagaku, 28 (1955) 527. 33 Gray R W and McCrum N G, J Polym Sci, A-2, 7 (1969) 1329. 34 Lee B L and Nielsen L E, J Polym Sci, Polym Phys Ed, 15 (1977) 683. 35 Nielsen L E , J Appl Polym Sci, 17 (1979) 1897. 36 Ashida M, Noguchi T and Mashimo S, J Appl Polym Sci, 29 (1984) 4107. 37 Becker G W and Oberst H, Kolloid Z , 152 (1957) 1 . 38 Woodward A E, Crissrnan J M and Sauer J A, J Polym Sci, 59 (1960) 23. 39 Ito E and Kobayashi Y, J Appl Polym Sci, 22 (1978) 1143. 40 Ito E and Kobayashi Y, J Appl Polym Sci, 25 (1980) 2145. 41 Willett P R, J Appl Polym Sci, 19 (1975) 2005. 42 Howard W H and Williams M L, Text Res J , 36 (1966) 691. 43 Mashimo S, Nakajirna M, Yamaguchi Y and Ashida M, Seni Gakkaishi, 41 (1985) T-225. 44 Mashimo S, Nakajima M, Yamaguchi Y and Ashida M, Seni Gakkaishi, 44 (1988) 606. 45 Mashimo S, Nakajima M, Noguchi T, Yamaguchi Y and Ashida M, Proceedings ofthe International Rubber Conference '87, Harrogate, (1987)50A; Rubber World, 200 (1987) 28. 46 Mashimo S, Nakajima M, Yamaguchi Y and Ashida M, Seni Gakkaishi, 41 (1985) T-448. 47 Mashirno S, Nakajima M, Yamaguchi Y and Ashida M, Seni Gakkaishi, 42 (1986) T-397. 48 Mashirno S, Nakajima M, Yamaguchi Y and Ashida M, Proceedings o f t h e International Rubber Conference '85, Kyoto (1985) 519.

6 Properties and processing of short metal fibre filled polymer composites D M BlGG

6 .I

Introduction

One of the most useful characteristics of polymers is that their properties can be modified to suit almost any need. Polymers can be rigid or flexible, transparent or opaque, insulating or conducting. Among the largest application areas for polymers are electrical and thermal insulators. In recent years there has been steady growth in the use of electrically conductive polymers. Conductive polymers can be either inherently conductive, or insulating polymers that have been filled with conductive materials. Inherently conductive polymers, such as polyanilines' and polypyrroles,2 are just beginning to find use in batteries, solar cells and superconductor applications. These polymers are specialty materials that cannot, as yet, be moulded or formed into large complex parts. Most electrically conducting polymers consist of familiar moulding resins that have been filled with conducting particles that form a conductive network within the polymer. Electrically conductive carbon black filled rubber has been used for many years to produce conductive membranes, buttons and gaskets. Carbon black filled polymers cannot provide sufficient conductivity to meet the requirements of the most predominant application for conductive polymers; that of providing electromagnetic interference (EMI) shielding3 A variety of worldwide governmental regulations, designed to protect the integrity of broadcast communications, has limited the permitted amount of electromagnetic radiation emitted from high frequency digital device^.^ Because there are significant advantages to housing such devices in plastic structures, methods are needed to provide shielding to a material that is ordinarily transparent to high frequency radiation. Electromagnetic theory shows that the only requirement for providing EM1 shielding is that the shield be capable of conducting electrical ~ u r r e n tIn . ~fact, the degree of conductivity required to provide various levels of shielding can be accurately calculated. For example, 40dB of attenuation can be provided by a material that has a resistivity of 0.01 R m . Figure 6.1 shows the relationship between shielding effectiveness and the resistivity of a material.6

Properties and processing of short metal fibre filled polymer composites 145

100

10

1

0.001

0.01

resistivity 6.1

0.1

1

(ohm m)

Relationship between shielding erectiveness and volume resistivity.

The key technical issues related to the use of filled plastics for EM1 shielding applications are: the type and form of the conductive filler used, the concentration required, the effect of moulding on mechanical and electrical properties, the ease of grounding the moulded structure to mating parts, and the long-term performance of the part. Some metals have active surfaces which can either oxidize or react with polymer^.^.^ The reactivity of metals primarily affects the minimum spacing between the particles required to ensure the conductive network. Long-term properties can be affected by reactions between the polymer and metal, and between permeable gases and the surfaces of the particles. Another application area in which metal filled polymers have been investigated is heat exchangers. In this application the low density, corrosion resistance and easy mouldability of polymers make them ideal candidates to replace metals. However, polymers have considerably lower thermal conductivities and upper use temperatures than metals, which places limits on the use of plastics in heat exchangers. Recent advances in the development of mouldable high temperature thermoplastics, having a heat distortion temperature above 250 "C, has opened up many potential application areas. These applications include automobile and truck radiators and heat exchangers in gas-fired high efficiency furnaces.' Efforts are still underway to enhance the thermal conductivity ofplastics for use in these applications. Higher plastic thermal conductivities will result in smaller, potentially less costly, heat exchangers than an all-metal unit. The addition of metal

I46

Short fibrepolymer composites

particles has been shown to increase the thermal conductivity significantly. l o Technical difficulties related to the effect of the metal particles on corrosion resistance, processability a n d mechanical properties are currently under investigation. A large body of literature generated over the past 15 years has addressed the problems related to the use of metal particulates in polymeric matrices. This chapter discusses the results ofthose developments with regard to the use ofmetal fibres and flakes to provide EM1 shielding and thermal conductivity enhancement in polymeric materials. Of particular interest are developments related to the primary metal fillers, stainless steel fibres and aluminium flakes.

6.2

Metal fillers

Aluminium, steel, iron, brass, copper, nickel and silver have been investigated in fibre o r flake form as conductive fillers in polymer composites. Each metal has certain attributes that make it a reasonable candidate for incorporation into polymers that are moulded into EM1 shielding structures. Because each metal also has a number of disadvantages, there is no ideal material for the application. Similarly, since heat exchangers often involve contact with corrosive chemicals, selection ofa metal filler must take the various material interactions into account. Aluminium has the lowest density, 2710kgm-3, and cost of the common metals. Aluminium is very conductive in bulk form, but since i t is covered with a surface layer of insulating aluminium oxide, small diameter aluminium powder is non-conductive. Aluminium has the lowest tensile strength of the common metals, 1 6 5 M P a , and is quite ductile. Aluminium will corrode rapidly in a n aggressive acidic environment. Steel is very low in cost, but since it has a very high density, 7800 k g m - 3 , its cost per unit volume is higher than that of aluminium. Stainless steels are relatively oxidation a n d corrosion resistant. Stainless steels are mechanically strong, but not as conductive as aluminium, copper o r silver. 316L stainless steel has a tensile strength of 550 MPa. Another low cost metal is iron, which has a similar density to steel, 7860kgm-3. Iron is more brittle and more conductive than steel, but less conductive than the other metals. Iron is more subject to oxidation a n d corrosion than steel, but considerably less so than aluminium and copper. Bras$ and copper are more costly than aluminium, iron o r steel, but cheaper than nickel or silver. Copper has a high density, 8820 kg m - and brass a density only slightly less at 8400 kg m-3. Both brass and copper have also been found to react chemically i n a degradative manner with many polymers. Nickel is more oxidation resistant than the other metals. It is only slightly more conductive than iron, a n d has been found to promote reactions in some polymers. Nickel has a density of 8900kgm-3. Silver is also subject to oxidation, has a very high density, 10 500 kg m - 3 , but is the most conductive ofthe candidate metals. Silver is also the most costly, a factor which has excluded it from consideration in most cases.

Properties and processing of short metal fibre filled polymer composites 147 In addition to consideration of basic bulk materials costs, there is an appreciable cost associated with converting metals into the desired fibre or flake form. Stainless steel fibres are produced by multiple wire drawing." In this procedure wire is drawn to a given diameter. Bundles of wires are then redrawn to a finer diameter, a procedure which is repeated until the desired diameter is achieved. This manufacturing process adds significantly to the cost of the fibres. Aluminium, copper, nickel and silver fibres can also be made by this technique. The resulting fibres are not smooth round fibres, as produced by the melt fibre spinning technique used to produce glass fibres. The fibres are produced continuously, then chopped to the desired length, usually after they have been coated with the polymer in which they will be embedded. Incorporation of the fibres into a polymer matrix is accomplished in a separate pultrusion process. Typical fibre diameters range between 2 and 20 p i , with 8 pm diameter fibres being the most commonly used for EM1 shielding applications. Types 301, 304 and 316 stainless steel are used to make the fine diameter fibres for addition to polymers. Short aluminium fibre and flakes have been made directly by a melt extraction process.'2 In this process a cooled spinning wheel is brought into contact with a reservoir of molten metal. Flake or fibre shaped details are machined or etched onto the surface of the wheel. The molten metal adheres to the wheel, solidifies, and is centrifugally ejected as distinct particles. Depending on the expense incurred in producing and replicating the cavities on the surface of the wheel, the resulting particles can have either a smooth rounded (never circular) surface, or a jagged irregular surface. Most of the flakes and fibrcs produced by this technique have been based on a low cost manufacturing approach, and the resulting particles have irregular surfaces. This process has other economic limitations. The cost of producing the cavities on the wheel increases significantly as the object particle size decreases. Moreover, even after the mould-wheel has been produced, the productivity decreases significantly when very small particles are produced. For that reason the particles produced by this technique which have been investigated represent a compromise between thc desired small particle size and low cost. Typical flake sizes are 1 mm wide x 4.25 mm long x 30 pm thick. The most frequently produced fibres are 1 mni long x 12.5 j.tm in diameter. Because the fibres are not smooth-walled cylinders they tend to aggregate during handling. The flakes do not aggregate, and are much easier to handle. This is the primary reason for the switch away from melt extracted fibres toward flakes. Low cost short fibres of aluminium, copper, steel and brass have also been produced by chatter machining.13 Fibres produced by this technique are 3 mm long and 60pm in diameter. Chatter machining does not produce smoothsurfaced cylindrical fibres. The important features of the spccific metal fibres and flakes that have been investigated are summarized in Table 6.1.

Thickness.

Drawing Melt extraction Melt extraction Chatter machining Chatter machining Chatter machining Chatter machining

Fibres Fibres Flakes Fibres Fibres Fibres Fibres

Stainless steel Aluminium Aluminium Aluminium Copper Brass Iron

a

Production method

Form

Metal 3-6 1.25 1.25 3 3 3 3

L (mm) 6-8 100 30" 60 60 60 60

D (Pm)

500-1000 12.5 42 50 50 50 50

LID

8020 2710 2710 2710 8900 8400 7900

Density (kgm-3) 74.0 2.8 2.8 2.8 1.7 7.0 9.8

Resistivity (10-80hmm)

Tuhle 6.1. Characteristics of metal fibres and flakes used in polymer composites

16.3 220 220 220 390 1 I7 71

Thermal conductivity K (Wm-' K-')

P

z

7

:

Y

2

V

7

5 0-

a

VI

Q)

Properties and processing of short metal fibre filled polymer composites 149 Table 6.2. Characteristics of polymers frequently used in conductive compounds

Polymer HDPE PC PP PPOjPS ABS Nylon 6,6 Pol ysulphone Poly(ether imide) PPS

Type Semi-crystalline Amorphous Semi-crystalline Amorphous Amorphous Semi-crystalline Amorphous Amorphous Semi-crystalline

Density (kgm-’)

Strength (MPa)

970 1,200 910 1,080 1,080 1,140 1,240 1,330 1,340

26 69 31 44 32 69 79 105 69

T, (‘C) - 20

150 0 108 105 50 190 217 85

T,,, (‘C) 132 -

I63 -

255 -

260

6.3 Polymers Almost any polymer can be used as the matrix in a metal particle filled electrically conductive structure. However, there are performance and economic requirements that have limited the number of materials that have actually been investigated. These materials include high density polyethylene (HDPE), polypropylene (PP), ABS, poly(pheny1ene oxide)/polystyrene alloy (PPO/PS), polycarbonate (PC), nylon 6,6 and poly(pheny1ene sulphide). Polymers used in thermal application areas include such corrosion resistant materials as poly(pheny1ene oxide)/polystyrene blend, poly(pheny1ene sulphide), polysulphone, poly(ether imide) and nylon 6,6. The characteristics of these polymers are shown in Table 6.2.

6.4

Electrical properties

Metals are excellent conductors, having resistivities in the order of l o p 8Rm, compared to the resistivities of polymers, which are in the order of 1013R m. In fine particulate form, however, the concentration of metal required to form a conductive composite is so large, approaching 40 v01%,’~that the material becomes difficult to mould, is quite dense, and lacks a satisfactory degree of mechanical integrity for most applications. Fortuitously, the electrical properties of a filled composite are a result of the development of a three dimensional network of conductive particles throughout the matrix. Such a network is formed at a critical concentration of filler particles. This critical concentration is reached when each particle makes contact with at least two neighbouring particle^.'^ The critical volume fraction is called the percolation threshold, Vf,,, an example of which is shown in Fig. 6.2.14This curve shows the change in resistivity of a phenolic composition which has been filled with randomly dispersed silver spheres. The percolation threshold is highly dependent on the geometry of the filler particles. Fibrous fillers have been shown to have relatively low percolation thresholds. l 5 The percolation threshold decreases with increasing fibre aspect ratio (length to diameter ratio). Such filler particles have a high probability of making ‘effective’ contact with neighbouring particles to form the required

Short fibre-polymer composites 1 O’O 10’ 108

1 o4

1o2

1 1o-2 i0-4

volume fraction of metal Effect of conductive filler concentration on resistivity ofcomposite filled with randomly dispersed silver spheres. (Source: Ref 14 with permission from the American Institute of Mining, Metallurgical, and Petroleum Engineers.)

network path for electrical conduction. Physical contact among particles is not necessary, but the spacing must be close enough (- 10 nm) to permit electron hopping across the gap between points of close contact; hence the use of the term ‘effective’ contact. l 6 The rate of hopping increases exponentially as the distance to be spanned decreases. The ability to make use of quantum phenomena to provide electron flow is the main reason why the presence of insulating surface coatings, such as are found on aluminium particles, significantly increases the percolation threshold of composites filled with coated particles. If the insulating surface layers are greater than the thickness through which electrons can hop, a conductive network will not develop, even if physical contact is made. This is the reason fine alumin’ium particles are not conductive. Numerous investigators have investigated the relationship between the fibre aspect ratio and percolation threshoid. Carmona et al proposed that the critical volume fraction was inversely proportional to the square of the aspect ratio.” From experimental data Bigg found that the relationship between aspect ratio and percolation threshold was inversely proportional to the 0.6 power of the aspect ratio. The experimental curve used to generate this relationship is shown in Fig. 6.3.18 This curve was developed from composites in which the fibres were randomly dispersed throughout the matrix. In most situations, however, the fibres and flakes become aligned in the flow direction during injection moulding (Fig, 6.4). This change in spatial arrangement affects the development of the network within the polymer. Fibre alignment generally increases the critical filler

Properties and processing of short metal fibre filled polymer composites I 5 I

1 o4

1o3

1o2

10

I

0.001

0.01

0.1

-

1

volume fraction of conductive filler 6.3 Effect of fibre aspect ratio on the percolation threshold in a randomly dispersed composite.

concentration required to induce conductivity, and makes the electrical behaviour of the composite anisotropic. This was recognized by Kortschot and Woodhams, who analysed the effect of both aspect ratio and degree of alignment on the electrical properties of composites filled with metal coated mica flakes. l 9 Unfortunately, their theoretical development was limited to two dimensional considerations of particle-particle interactions, which even in the simplest case of spherical particles are invalid. Wang and Ogale developed a theory to predict the percolation threshold of polymers filled with short fibrous conductors.20 This theory accommodates the specification of a variable ‘effective’contact distance beyond the actual physical dimensions of the fibres. As with all theories developed to predict the percolation threshold, this one is based on Monte Carlo statistical computational techniques. Figure 6.5 provides a comparison of the theoretical predictions developed by Wang and Ogale with data generated by two sets of investigators. The theory provides a reasonable starting point for determining the percolation threshold of fibrous or flake filled composites. At the present time extension of the theory to fibres with aspect ratios greater than 25: 1 requires excessive computer time. De Bondt et al also provided a three dimensional percolation model to predict the conductive network threshold based on considerations of particle shape, particle size and particle orientation.2’ They commented on the length of

I52

Short fibre-polymer composites

6.4

Cross-section of iiijcction inouldcd part showing parallel alignment of a I uni i n i urn flakes.

computer time required to complete an analysis, and concluded their evaluation of the effect of aspect ratio at an L / D ratio of 10: 1. Their results are in good agreement with published experimental data for isotropically distributed particles having aspect ratios between 1 and 10. Metal fibres are ductile, a complicating factor that must also be taken into account. I n ;I viscous flow field, subjected to a high shear gradient, metal fibres will bend, and n o t straighten upon cessation of shear. Folkes speculated that the multidimensional shear field present in most moulding machines bends ductile fibres into helical coils, which increases each fibre's probability of coming into c r i t ica I pro x i m i t y with its neigh bo u r s. T o ;icco mm od a t e this distortion of m e t al fibres, t o l k e s hypothesized that the percolation threshold can be predicted by the following eclu;~tion:

'

c;,c

= xr.2S/100D2/l

C6.11

i n which I' is the radius of tlie fibres, S is tlie length of the average 'coil', I1 is the diameter of tlie cross-section swept out by tlie helix, and h is tlic thickness of tlie moulded plaque. A n interesting concliisioti of Folkes' work is that long conductive fibres induce greater conductivity in polymeric matrices by virtue of their propensity t o distort than they would havedone liad they been rigid fibres of the same length. The formation of helical coils during flow provides for an increase in probable contact sites, relative to a rigid fibre. This is easily demonstrated. Assume ;I fibre radius of6 pin, ;I part thickness of 3 mm and a rigid

Properties and processing of short metal fibre filled polymer composites I53

1

0.1

0.01

I

I

1

1

1

1

1

I

1

I

rn ‘

I

,

100

10

aspect ratio 6.5 A comparison of Wang and Ogale’s theoretical predictions (0, 0) of the percolation threshold with experimental data from Ref 6 ( 0 )and 19 (B). (Source: Ref 20 with permission from Elsevier Applied Science.)

length of 6 mm. Vf,cfor a composite in which these fibres are maintained rigid is 0.0157. If during processing this fibre is distorted to form a coil in which D becomes 21 pm, and S is reduced to 5mm, Vf,, is reduced to 0.0044. This relationship is purely empirical, and breaks down when rigid fibres or fillers with other shapes are considered. Once the percolation threshold has been found, whether theoretically or experimentally, the relationship between composite resistivity, pc, and volume fraction of filler, Vf, in the neighbourhood of the percolation threshold can be described by the simple r e l a t i ~ n s h i p : ~ ~ Pc

=

KC( Vf

-

l’f,c)/Vf,cl

-‘

C6.21

where t is a power law coefficient between 2 and 3.5 for many compositions, and K is a coefficient approximately equal to 0.15. This equation does not cover the behaviour of the composite prior to network formation, or in the region beyond the percolation threshold. I t is importimt because the relationship in the percolation region is of considerable interest in the development of positive temperature coefficient, PTC, materials. To date, PTC composites are based exclusively on conductive carbon black. McLachlan developed an equation which describes the entire ‘Z’ shaped

Short fibre-polymer composites 10’0 108 108 10‘ 102

1 10-2

10-4 0

0.1

0.2

0.3

volume fraction o f conductive filler Resistivity versus volume fraction for aluminium flake filled composites, as predicted from eq C6.31.

resistivity versus volume fraction of filler curve that covers the transition from insulator to conductor.24This relationship makes use of the asymptotic values of resistivity as Vf approaches 0 and the maximum packing fraction, Vr.n,ax,as well as knowledge of Vr,c.This is basically a curve fitting approach that relies on the fact that the shape of the resistivity-concentration curve is quite similar for most composites filled with conductive fillers. The equation developed by McLachlan is:

where p p is the resistivity of the polymer, p, is the resistivity of the composite a t V,, pr is the resistivity of the composite at Vr,maxand s is an exponent that must be determined. The utility of this equation depends on the ability to determine the percolation threshold. It has significant value in that it allows predictions of the composite resistivity in the concentration region above the critical threshold, the region that is important in the production of conductive EM1 shielding materials. Figure 6.6 shows the predicted relationship for a n aluminium fibre filled composite having a percolation threshold of 0.1 3, with a value of s equal to 1.O. This curve provides a reasonable approximation to the curve experimentally determined for injection moulded aluminium flake filled composites, shown in Fig. 6.7. The most significant problem with equation [6.3] is that it is not yet possible to

Properties and processing of short metal fibre filled polymer composites 155

E E

r 0

)r

.-w .-ww

.2 v)

u)

6.7 flake filled composites. Filled symbols: compression moulded parts, open symbols: injection moulded parts, polypropylene, A polycarbonate, 0 ABS.

predict pr. If the conductive filler consists ofcontinuous, unidirectionally oriented fibres, the bulk resistivity of the composite in the direction of the fibres is simply the resistivity of the fibrous material divided by the volume fraction of fibres. Because the fibres in a moulded structure have some degree of random orientation, the resistivity of the composite will be somewhat higher than that ofa unidirectionally oriented structure. In a fibre reinforced composite the modulus decreases by approximately a factor of three when the fibres are randomized within a plane. By analogy, the relative increase in resistivity of a conductive composite in which the conductive fibres are randomly oriented within a plane should be of the order of a factor of three. Even this magnitude of increase is exceeded by a considerable margin in metal filled composites. The reason for this disparity is that factors other than geometrical considerations must be considered. There is a significant level of electrical resistance offered at each particle-particle contact point that is not accounted for by geometrical considerations. I t was discussed earlier that these contact points are not always points of actual physical contact. In many, if not most, cases the contact points are simply locations close enough to allow electron hopping to occur. Because the rate of electron hopping is related in an exponential manner to the distance to be hopped, there is a variation in the ability of electrons to hop at each contact point. Since the ability to hop decreases with distance in an exponential manner, the contact point resistivity is much greater than that

I56

Short fibre-polymer composites

assumed in the simple models. Therefore, the overall bulk resistivity is much higher than that predicted from geometrical considerations alone.

6.5

Thermal properties

The effect of filler particles on the thermal conductivity of a polymer is not a network phenomenon, as electrical conductivity is, but a bulk effect. The thermal conductivity of the composite increases monotonically as the filler concentration increases. The rate at which the thermal conductivity increases depends on the thermal conductivity of the filler particle, its shape, and degree of dispersion within the matrix. There are many theoretical developments describing the effect of various fillers on the thermal conductivity of a heterogeneous composition. One versatile and useful relationship is that developed by N i e l ~ e n Nielsen's .~~ model, equation C6.41, has as key parameters the maximum packing fraction Vf.maxand a factor A , which is related to the geometry of the filler particle. A has a value of 1.5 for spherical filler particles; values for several common filler types are listed in Table 6.3. kJk

=

1 +ABV,

1

-

B$

vs

c6.41

where B=

kflk - 1 k,lk A

+

~6.51

and

Figure 6.8 shows the effect of aluminium fibre and flake concentrations on the thermal conductivity of a number of polymeric composites. The curves in this graph are based on the model developed by Nielsen. There is good agreement between the model and experimental results. This model has also been found to provide reasonable estimates of the thermal conductivity of other composites. ' O,*

6.6

Mechanical properties

There are several good reasons to reduce the concentration of metallic fillers used to produce conductive composites. These include considerations of cost, density, processing difficulty and mechanical property performance. Of these considerations, the effect of metal fillers on mechanical properties has been the most significant concern. Polymers exhibit poor adhesion to metals, which makes the resulting composites susceptible to reductions in tensile strength and most significantly losses in impact strength.

Properties and processing of short metal fibre filled polymer composites I57 Table 6.3. Parameters A and V,,,,, for metal fibres and flakes Filler Spherical particles Irregular particles Aluminium flakes Steel fibres

Dispersion

LID

A

Random Random Random Random

1.0 1.0 10.0 500.0

1.5 '-2.5 -5 -500

~~f,,,,

0.637 0.64 0.33 0.01

10

1

0

0.1

0.2

0.3

volume fraction of metal fibres Effect of aluminium flake concentration on thermal conductivity of polymer composites: (0) Ref 26; ( 0 )Ref 27; ( 0 ,W ) Ref 28; - l/d = 12.5, K,JK, = 1000, Vl.mar= 0.33; ....... l/d = 12, K,/K, = 212, V,.,,, = 0.52; - - l/d~= 10, ~ K , / K , = 1000, V,,,,, = 0.52.

6.8

The tensile strength of short fibre filled composites is dependent upon both the degree of adhesion between the polymer and filler and the aspect ratio of the fibres. For each polymer-fibre combination there is a critical aspect ratio for maximum reinforcement, that critical aspect ratio being related to the adhesion between the polymer and fibre. The critical aspect ratio for maximum reinforcement is ('l4crit

= 'f,crit/'T7

c6.71

where is the tensile strength of the fibre and T? is the shear strength between the polymer and fibre (see Chapter 1, section 1.4.1). In general, polymers do not

I58

Short fibrepolymer composites

Tahle 6.4. Effect of A-1 100 silane coupling agent on the tensile strength of aluminium fibre tilled composites Matrix

v,

Coupling agent

Nylon 6,6 Nylon 6,6 Nylon 6,6 Pol ycarbonate Polycarbonate Polycarbonate

0.0 0.2

No No Yes No No Yes

0.2

0.0 0.2 0.2

=<,c,,1

(MPa)

68.9 37.1 55.1 69.0 35.6 39.0

readily wet and adhere to metals, hence the adhesive shear strength is quite low. This results in very high values for the critical fibre aspect ratio. In order to obtain the maximum degree of reinforcement in a composite reinforced with randomly dispersed fibres, the fibres in the composite should have an aspect ratio ten times as high as the critical aspect ratio.30 Therefore, if the critical fibre ratio is equal to or higher than the actual fibre aspect ratio, no reinforcement will occur, and the tensile strength of the composite will be less than that of the pure material. Some success in promoting adhesion has been obtained by using coupling agents and chemically modified polymers. Table 6.4 shows the effect of silane coupling agents on the tensile strengths of aluminium fibre filled nylon 6,6 and polycarbonate composite^.^ Figure 6.9 shows the effect of maleic anhydride modified polypropylene on the tensile strengths of metal filled composite^.^^ Polyvinyl alcohol, PVA, surface coatings are often used for sizings on stainless steel fibres. PVA coatings have also been shown to be effective in promoting adhesion between polypropylene and an amorphous metal fibre c o m p ~ s i t i o n . ~ ~ O n a more pragmatic basis, not much mechanical property reinforcement should be expected from metal fibre filled composites. Steel and aluminium have tensile strengths that are far less than those of common reinforcing fibres, glass and carbon fibres. Glass fibres have a tensile strength of 2.4 GPa. Carbon fibres, which have numerous grades, also have tensile strengths as high as 2.4 GPa, and are even used to reinforce metals. A common method of representing the tensile strength of fibre filled composites is the modified r ule- of -mi~ tur es:~ ~

’

where: crc,crit= tensile strength of the composite, = tensile stress in the polymer matrix at failure, volume fraction of the polymer matrix in the composite, = tensile strength of the fibre, Vf = volume fraction of the fibre in the composite, E~ = fibre orientation factor, and E~ = reinforcement efficiency.

V,

E~

=

is 1 for unidirectionally oriented fibre, 0.375 for randomly oriented fibres in a

Properties and processing of short metal fibre filled polymer composites 159

1.2

I

I

1

1

t

1

1.o

0.8 0.6 Q,

.->

0.4

CI

a Q,

L-

0.2

'

01 0

I

I

I

0.1

0.2

0.3

I

0.4

volume fraction of filler 6.9 Effect of chemically modified polypropylene on the tensile strength of aluminium fibre filled composites: ( 0 )chemically modified polypropylene composites; ( 0)unmodified homopolymer polypropylene.

planar orientation, and 0.16 for random three-dimensionally oriented fibres. E depends on the degree of polymer-fibre adhesion, and, consequently, the critical aspect ratio. For strongly adhering glass fibre-polymer combinations, the critical aspect ratio is between 15: 1 and 30: 1. While this is less than the aspect ratio of most steel fibres and comparable to the aspect ratio of aluminium fibres and flakes, neither of these fillers provide much tensile strength enhancement. Stainless steel fibres are used in concentrations of approximately 1 ~ 0 1 %At . this concentration, the calculated tensile strength of the composite, assuming adequate polymer-fibre adhesion, a random in-plane fibre distribution, and a polymer tensile strength of 70 MPa, is 71 MPa. The concentration of aluminium flakes used in many electrically conductive formulations is in the order of 20 ~01%.The calculated tensile strength of an aluminium flake filled composite is 65 MPa, a slight reduction in strength compared with the unfilled polymer. In Fig. 6.9 the effect of filler concentration and degree of polymer-filler adhesion on the tensile strength of metal fibre filled composites was shown. These curves follow the relationship developed for filled compounds, rather than those developed for reinforced materials:

I60

Short fibrepolymer composites Table 6.5. Coefficients for equation [ 4 . 9 ] from curves in Fig. 6.9

Curve

(1

Upper curve Lower curve

1.50 1.21

b

c

e

0.67 0.67

0.61 0.00

0.33 -

yield

drawing

b

I I

I I

strain: 6.10

AL/L

Stress-strain curve of typical ductile polymer. (F/A, = engineering

stress.)

c6.91 In this equation, CI is a constant related to stress concentrations in the composite, and h is a constant related to the geometry of the filler. T h e parameter h often has a value of 2/3. c and e a r e coefficients related to the small degree of reinforcement many filled systems exhibit a t low filler concentrations. c often has a value of 0 for many compounds filled with spherical o r irregularly shaped particles. Table 6.5 shows the values of N , h, c and e for the two curves shown in Fig. 6.9. The tensile properties of most polymers are satisfactory for most applications, so the lack of strength enhancement is not a serious detriment. A very serious shortcoming is the reduction in impact strength. Impact strength is difficult to measure and predict in a consistent manner. The impact strength of a plastic material is very dependent on the technique used to measure it, the size of the sample, the presence o r absence of a notch, and the orientation of the sample.

Properties and processing of short metal fibre filled polymer composites I6 I

700 CI

b~

600

Y

A

m

500

Q) L

400

\

0

milled glass fibres Al fibres (L/D = 12.51

0

CI

.-

300

0

CI

m 200 c

-0 Q)

100 0

0

0.1

0.2

0.3

0.4

0.5

volume fraction of filler 6.11

Effect of filler concentration on the strain-at-break for polypropylene composites.

Data generated in one impact test cannot be correlated with data collected in another impact test. The impact response of a material is the amount of energy it can absorb during an impact. Recognizing that polymers are non-linear viscoelastic materials, this means that the impact behaviour is also rate-dependent. In very basic terms the energy absorbing characteristics of a material can be represented as the area under a stress-strain curve. A typical stress-strain curve for a ductile, unfilled polymer is shown in Fig. 6.10. As particulates are added to such a polymer, the ultimate strain is reduced (see Fig. 6.1 1) regardless ofwhether the strength of the composite increases or not. In reinforced composites, even though the ultimate strain is reduced, the strength and modulus often increase enough to more than compensate for the reduction in ultimate strain. Therefore, in many reinforced composites the area under the stress-strain curve will increase despite the reduction in ultimate strain. If the tensile strength of the filled composite does not increase, then the area under the stress-strain curve must decrease. For example, consider a polymer that has a tensile yield and breaking strength of 70MPa, a yield strain of 0.1 and an ultimate strain of 0.6. Upon adding sufficient metal fibres or flakes to produce the desired amount of conductivity, the ultimate strain drops to 0.1, but the tensile strength remains at 70 MPa. The approximate area under the stress-strain curve for the polymer is 38.5 MPa. The area under the stress-strain curve for the filled polymer is reduced

I62

1 .o

r w c

0.8

a L

CI v)

0.6

0

a

Q

.-

0.4

a

>

I-

w

-ma

0.2

L

0 0

0.1

0.2

0.3

0.4

volume fraction of filler 6.12

Effect of metal filler concentration on the impact strength of poly-

propylene composites: ( 0 )aluminium flake filled composites; ( 0 )silicon carbide particulate filled composites.

to 3.5MPa, which is less than 10% of the energy absorbing capability of the unfilled polymer. Figure 6.12 shows the effect of metal filler concentration on the impact properties of conductive composite^.^^

6.7 Processing There are two aspects to processing metal fibre and flake filled composites. The first consideration is mixing the filler into the polymer with as little damage to the filler particles as possible. The second consideration is moulding the conductive compound into the desired shape, also with as little damage to the particles as possible. Each of these process operations is considered separately. 6.7. I

Compounding

Fillers can be added to polymers by a wide variety of techniques. These include simple dry blending, melt blending by either continuous or batch methods, extrusion coating and pultrusion. All of these techniques have been investigated for metal filled conductive composites. Regardless of the application it is important to minimize the damage done to the high aspect ratio particles. Simple

Properties and processing of short metal fibre filled polymer composites I63

6.13 Photograph of steel fibres removed from injection moulded cornposites showing reduction in length and coiling.

dry blending does not provide adequate dispersion because of the mismatch in densities and particle geometries between common metal fillers and most polymers. Compounding of discrete particulates, such as aluminium flakes, is best accomplished by adding the filler to the polymer melt in a low shear, distributive mixing extruder. Continuous fillers, such as the stainless steel fibres, can be added to the polymer without any size reduction by either extrusion coating o r pultrusion. Pultrusion provides complete wetting of the fibres by the polymer, a factor which reduces the mixing required in the moulding machine to produce a homogeneous dispersion i n the part. Wetted filler particles are not as easily damaged during the melting process in the moulding machine as the unwetted fibres that come from the extrusion coating process. 6.7.2 Moulding In commercial practice the most common forming process for thermoplastic polymers is injection moulding. Unfortunately, many of the studies of properties versus filler concentration relied on compression moulding, which does not provide the same degree of shear to the fibres as injection moulding. Folkes and co-workers showed that 6 mm long stainless steel fibres, having a diameter of 6.5 pm, were reduced in size during single screw extrusion compounding and injection moulding to an average length of approximately 0.5 mm, a significant reduction in length.22 Figure 6.13 shows both the reduction in size of stainless

I64

Short fibrepolymer composites AS MANUFACTURED

AFTER

PROCESSING

ALUMINUM FLAKES 6.14

Photograph of aluminium flakes removed from injection moulded composites showing tearing and folding of flakes.

steel fibres during injection moulding and the coiling that occurs. Because of the ductility ofstainless steel fibres the effective fibre length is shorter than the actual length of the fibres. Most fibres are kinked, bent, curved or folded. Aluminium flakes have also been shown to be significantly damaged during melt processing. Figure 6.7 showed the effect of the shear forces on the electrical properties of aluminium flake filled composites. Composites produced by compression moulding exhibited a percolation threshold of 0.075. Injection moulding increased the percolation threshold to 0.13, almost twice as high. All of the damagedone to the flakes was a result of the injection moulding process, since dry blending was used to produce the mixture fed to the moulding machine. Figure 6.14 shows the damage inflicted upon the flakes during injection moulding. There are tears in the flakes and folds, both of which effectively reduce the aspect ratio of the flakes. Recent advances in process technology have resulted in fabrication techniques that can produce thin-walled, complex-shaped parts that do not degrade the filler particles to the same degree as the injection moulding process. These process techniques also maintain the homogeneity of the dispersion to a greater degree than injection moulding. The first process is compression blow moulding;35 the second is fast compression moulding of sheet c o m p o ~ i t e s . ~ ~

Properties and processing of short metal fibre filled polymer composites I65

Compression blow moulding has been developed to produce thin-walled, hollow structures from filled engineering thermoplastics. In this process a round parison is extruded and pinched at both ends, one end at a time. Prior to pinching the second end the parison is pressurized. The pressure in the parison is not high enough to inflate the parison, since the filled polymer does not have enough melt strength to expand to a significant degree. The pressure in the closed parison prevents it from collapsing during compression of the soft parison by the closing mould. Since there is no flow through thin channels, there is little shear damage beyond that done in the extruder and no shear induced unmixing of the filler from the polymer. Another positive aspect of this process is that there are no weld lines. Weld lines can be significant sources of radiation leakage in EM1 shielding components. Compression blow moulding is being used to produce stainless steel filled shielding components. Thermoplastic sheet composites also produce parts without weld lines, but demixing will occur if the sheet is moulded to very high extension ratio^.^' Typically, the sheet is not extended beyond a thickness reduction of 2: 1. Thermoplastic sheet is produced by either melt impregnation or slurry deposition from an aqueous dispersion. Both processes produce composites with fibres that can be longer than 3mm, and are more typically longer than 12mm. The compression moulding process used to convert the sheet into parts does not impart much shear, so the fibre length is maintained. One of the primary attributes of sheet composites is very high impact strengths, as a result of maintenance of fibre length.38 This technique has been investigated to produce EM1 shielding material^,^' but has not been commercialized.

6.8

Effect of environment o n properties

6.8.I

Electrical properties

When EM1 shielding requirements first became necessary, electrically conductive coatings were applied to the interior of moulded plastic structures to provide shielding. At the time, concern was expressed that such coatings would delaminate from the moulded structure during the repeated thermal cycling between the -20°C and +80"C range that electronic components are required to withstand.40 It was the anticipated stability of filled conductive compounds that provided early encouragement for the development of filled compositions. Improvements in conductive coatings have prevented delamination from becoming a serious problem, and nickel flake filled acrylic paints have become the standard material for providing EM1 shielding to moulded plastic structures. Several studies have examined the effect of thermal cycling on the electrical performance of filled composites, and in many cases the anticipated electrical stability was not Osawa and Kobayashi showed that copper and aluminium fibre filled polyethylene composites rapidly lost their conductivity when heated at 80 " C 8 Brass and iron lost conductivity, as well, but to a much lower degree. The loss in conductivity in the aluminium and copper fibre filled

Short fibre-polymer composites

I66

composites was attributed to oxidation of the metal at elevated temperatures. Bigg showed that certain compounds filled with aluminium flake, nickel flake and stainless steel fibres lost substantial conductivity after five cycles between - 20 "C and 80 0C.41The compounds that lost conductivity were those that had lower temperature polymer matrices, ABS, poly(pheny1ene oxide)/polystyrene blend and PVC. It was hypothesized that these materials had sufficient molecular mobility that the conductive network was reduced under the repeated expansion and contraction cycles. Because the metal fillers and polymers have different coefficients of thermal expansion there was a difference in the degree and manner in which the two phases moved during thermal cycling. Similar compounds based on higher temperature polymers, such as poly(pheny1ene sulphide), nylon 6,6 and SMC, did not lose much conductivity after thermal cycling. Not surprisingly, these materials have less molecular mobility and lower coefficients of thermal expansion in the temperature range investigated. Mobius also reported that stainless steel filled thermoplastic composites irreversibly lost electric conductivity during thermal cycling.42 Another environmental stress that some conductive composites can experience during service is from harsh chemicals. This is a significant problem for electronic control devices in chemical processing plants and under-bonnet automobile applications. Bigg showed that aluminium flake and stainless steel fibre filled compounds lost significant conductivity in both 12 week controlled laboratory exposure in a detergent solution, and year long outdoor exposures at widely scattered chemical process plants when the matrix polymer was a low temperature material such as ABS, poly(pheny1ene oxide)/polystyrene blend or PVC.43 Compounds based on such high temperature polymers as poly(pheny1ene sulphide) and SMC did not exhibit any reduction in electrical conductivity.

+

6.8.2 Thermol properties The bulk thermal properties of a metal filled polymer will not be reduced by chemical attack, but the long-term structural performance of the material may be affected. There are few studies on the effect of environmental conditions on the performance of metal filled polymers in heat transfer applications. Bigg et al studied the effect of an aluminium flake filled poly(ethy1ene terephthalate) composite as the condensing heat exchanger of a natural gas furnace.' Extruded tubes of this composite were subjected to 10 000 cycles of 11 min with the furnace on and 4min with it off. In these experiments the incoming flue gas peaked at 190 "C. The flue gas was spiked with trichloromonofluoromethane to produce a condensate with 26 ppm chloride ion and 5 ppm fluoride ion, which is typical of the ion content in the condensate of residential gas furnaces. The acidic condensate severely attacked exposed aluminium flakes in the composite tubes.

References 1 Epstein A J and MacDiarmid A G , S P E A N T E C , 37 (1991) 755. 2 Kuhn H H, Worrell W C G and Chen C S, S P E A N T E C , 37 (1991) 76.

Properties and processing of short metal fibre filled polymer composites I67

3 Bigg D M, J Rheol, 28 (1984) 501. 4 FCC Rules and Regulations, Vol. 11, Part 15, Docket No. 20780 (July 1981). 5 White D R J, E M C I / E M C Hundbook Series, Vol. 4, Don White Consultants, Germantown, M D (1984). 6 Bigg D M and Stutz D E, Polym Compos, 4 (1983) 40. 7 Kelleher P G and Miner R J, S P E A N T E C , 30 (1984) 385. 8 Osawa Z and Kobayashi K, J Muter Sci, 22 (1987) 4381. 9 Bigg D M, Stickford G E and Talbert W E , Polym Eng Sci, 29 (1989) 1111. 10 Bigg D M, Polym Compos, 7 (1986) 125. 11 Tolokan R P, S P E A N T E C , 30 (1984) 699. 12 Maringer R E and Mobley C E , J Vuc Sci Technol, 11 (1974) 1067. 13 Nakagawa T, Suzuki K and Koyama H, Metul, 50 (1980) 5. 14 Gurland J, Trans Metal SOCA I M E , 236 (1966) 642. 15 Bigg D M, Polym Eng Sci, 19 (1979) 1188. 16 Bridge B, Folkes M J and Jahankhani H, J Muter Sci, 23 (1988) 1955. 17 Carmona F, Barreau F, Delhaes P and Canet R, J Phys Lett, 41 (1980) L531. 18 Bigg D M, 'Metallfilled Polymers - Properties and Applicutions', Ed S K Bhattacharya, Marcel Dekker, New York (1986) 202. 19 Kortschot M T and Woodhams R T , Polym Compos, 9 (1988) 60. 20 Wang S F and Ogale A A, Compos Sci Technol, 46 (1993) 93. 21 De Bondt S, Froyen L and Deruyttere A, J Muter Sci, 27 (1992) 1983. 22 Bridge B, Folkes M J and Jahankhani H, J Muter Sci, 24 (1989) 1479. 23 Chen I G and Johnson W B, J Muter Sci, 26 (1991) 1565. 24 McLaclan D S, J Phys, C20 (1987) 865. 25 Nielsen L E , Ind Eng Chem, 14 (1974) 17. 26 Hamilton R L and Crosser 0 K, Ind Eny Chem, 1 (1962) 187. 27 Garrett K W and Rosenberg H M, J Phys D: Appl Phys, 7 (1974) 1247. 28 Bigg D M, Composites, 10 (1979) 85. 29 Progelhof R C, Throne J L and Ruersch R R, Polym Eng Sci, 16 (1976) 615. 30 Folkes M J, 'Short Fibre Reinforced Thermoplastics', Research Studies Press, Letchworth, UK (1982). 31 Bigg D M, J Ind Fabrics, 2(3) (1984) 4. 32 Bigg D M, Polym Compos, 8 (1987) 115. 33 Sidhu A S and Varin R A, Polym Compos, 14 (1993) 277. 34 Bigg D M, Hiscock D F, Preston J R and Bradbury E J, Polym Compos, 9 (1988) 222. 35 Higgins J, S P E A N T E C , 34 (1988) 763. 36 Bigg D M and Preston J R, Polym Compos, 10 (1989) 261. 37 Hojo H, Kim E G and Tamakawa K, Int Polym Proc, 1 (1987) 60. 38 Bigg D M, Polym Eng Sci, 32 (1992) 287. 39 Yats L D and Edens M W, Tuppi J , Aug (1988) 81. 40 Regan J, Polym Plust Technol Eng, 18 (1982) 47. 41 Bigg D M, Polym Compos, 7 (1986) 69. 42 Mobius K H, Kunststoffe, 78 (1988) 53. 43 Bigg D M, Polym Compos, 8 (1987) 1.

Electrically conductive rubber and plastic composites with carbon particles or conductive fibres P B J A N A , A K M A L L I C K A N D S K DE

7. I

Introduction

Elastomers and plastics contain very low concentrations of free charge carriers. They are electrically non-conductive and transparent to electromagnetic radiation. T ~ L IifSthey are used as enclosures for electronic equipment they can neither shield it from outside radiation nor prevent the escape of radiation from the components. As such they are unable to function as electrostatic charge dissipating materials and semiconducting materials used in handling sensitive electronic devices. To provide conductivity a n d shielding from electromagnetic interference (EMI) requires incorporation of fillers of high intrinsic conductivity such as particulate carbon blacks, carbon and graphite fibres, metal particulates, and metal fibres to the polymer matrix. A critical concentration of filler, beyond which the polymer composite changes from an insulator to a conductor, is referred to as the percolation threshold. At this point a continuous network of filler particles is formed throughout the polymer matrix. This allows the movement of charge carriers present in the filler throughout the polymer matrix, thus making the composite electrically conductive. Electrical conduction of materials may occur through thc movement of charge carriers which follows the equation:' (T

=

erlp

~7.11

where (T is the conductivity, r the charge, ti the carrier concentration and p the mobility of the carriers. In electroconductive polymer composites the formation of the conductive network, the nature of the charge carriers and their concentration and mobility depend on filler concentration, filler geonictry (size and shape) and their distribution, and processing conditions.' ' The processing conditions include the preparative techniques, degree of vulcanization and mould pressure.

7.2 Percolation phenomena in conductive polymer composites The most convincing theoretical approach to studying the percolation behaviour

Electrically conductive rubber and plastic composites

I69

of a composite containing a random mixture of conducting particles in an insulating matrix is in terms of percolation theory, as developed by Kirkpatrick.’ According to the theory, a composite is regarded as a lattice of conductive ‘sites’ joined by resistive ‘bonds’ and the conductivity (u) of the composite above the percolation threshold follows a power law dependence of the form of c7.21 where p , is the critical or threshold probability of formation of a conducting network, p is the probability of finding the conducting phase which is equivalent to the volume fraction of conducting phase above the critical concentration, o o is the pre-factor and t is a critical exponent. According to the most recent studies6 the value o f t is around 2.0 for three dimensional composite materials. A t the critical or threshold concentration a sharp transition between insulating and conducting states has been observed. Several researchers observed a power law dependence of conductance on filler concentration in the polymer composi t e ~ . ’ -Alternative ~ approaches to the interpretation of percolation behaviour of electrical conduction have also been made. Percolation in randomly oriented short fibre filled epoxy resin composites showed a significantly larger value of critical exponent (t) than the universal value (t = 2.O).I4 Balberg and Bozowski observed good agreement between the theoretical prediction of the critical exponent and the value determined experimentally for the percolation in a composite of stick-like conducting particles.” Percolation phenomena of the carbon fibre filled rubber composites have been studied at different fibre aspect ratios.I5 It was observed that the critical exponent ( t ) is around 2.0, which is independent of the fibre aspect ratio. The critical concentration or percolation threshold at which a continuous conductive network of filler particles is formed in the polymer matrix depends on various factors.’ Processing conditions greatly affect the threshold value.“ An increase in dispersion time tends to break down most of the carbon aggregates and to coat each carbon particle with a layer of polymer, which results in a high threshold value. Dependence of the threshold concentration of the conducting phase on the surface tension between the polymer and carbon particles has been reported.’ i 9 1 8 The volume fraction of fillers required to reach the percolation threshold decreased when highly structured and extended carbon black particles were used to make conductive polymer composite^.'^ The effect of structure on the percolation threshold was also investigated by Janzen,20 who emphasized the mean number of contacts between the filler particles, rather than the probability of ‘bond’ or ‘site’ occupancy as considered in percolation theory. The particle size ratio of the filler and the polymer also affects the threshold concentration. Filler particles with larger surface area than polymer particles would lower the percolation threshold.21*22 It has also been observed that a composite containing a chain-like structure of fine carbon black particles exhibits a lower percolation threshold as compared with one containing isolated clusters of carbon black.23 Although the experimental results were always concerned with isotropic percolation, anisotropic percolation has also been investigated in conductive

1 o8

1

I

I

I

” -a

106

-

1

.

-

I I

104 -

,,

-

I I !

I

102 -

,, , , ,

-

t

1 -

1 8

-

Q.. 10-2 -

--Q---

- - a -- - -0 - - - - - - - - -

Q

(b) I

10-4

I

I

I

Effects of CF concentration and fibre aspect ratio ( L / D ) on volume resistivity: (a) composites prepared by mill-mixing method (LID = 25); (b) composites prepared by cement-mixing method (LID-100).(Source: Rcf26 with permission of the publisher, American Chemical Society, Inc.) 7.1

polymer composites. Balberg et al have reported that the resistivity dependence on the me1t flow behaviour of carbon black-PVC composites can be explained by an anisotropic percolation The conductivity behaviour of randomly distributed carbon fibre-polymer composites near the threshold has been studied by Jana and c o - ~ o r k e r s . ’ ~It. was ~ ~ , found ~ ~ that the composite containing a high fibre aspect ratio, prepared by the ‘cement’ mixing method, needs a lower critical concentration of fibre than the composite with low fibre aspect ratio prepared by the ‘mill’ mixing method (Fig. 7.1).26 Randomly distributed filler particles in a polymer matrix are usually described by the percolation theory. It has been reported that the filler particles can also be distributed in the polymer matrix in a segregated fashion, rather than in a random fashion, leading to the formation of a conductive n e t ~ o r k . ~The ’ , ~percolation ~ theory cannot be applied to the segregated network system. The segregated network is achieved through the tendency of the fine filler particles to adhere to the surface of much larger particles of polymer. More recently Yacubowicz et al reported that segregated distribution of carbon black in polymer matrix lowers the percolation threshold or critical concentration value significantly, as compared with the threshold value of a randomly distributed system.”

7.3

Mechanism of electrical conduction

Depending upon the types of filler and their concentration, various mechanisms of conduction have been proposed which include tunnelling of electrons, electron

Electrically conductive rubber and plastic composites

171

hopping, ionic transport, field emission, band type conduction and simple inter-aggregate conduction. Polley and Boonstra proposed that conduction in carbon black composites is achieved by means ofelectrons jumping across the gap between one carbon black aggregate to another, with an exponential dependence on gap width.” These composites show non-ohmic behaviour of electrical conduction and are believed to follow a tunnelling mechanism of c o n d ~ c t i o n1,32 . ~ The tunnelling of electrons through the gap is regarded as a special case of internal field emission. A review by Medalia on electrical conduction in carbon black composites summarizes the factors controlling the tunnelling of electrons in these composite^.^ Tunnelling is a quantum mechanical process in which the wave function of the electron is not confined entirely within a potential barrier, but has a small tail extending beyond the barrier.33 Owing to the presence of the tail, a very small fraction of electrons penetrate the barrier and reach the next carbon black aggregate and thus the conductivity of the composite is achieved. and M ~ t have t ~reported ~ both band type conduction and hopping of electrons. Band type conduction occurs when the charge carriers (electrons or holes) are transferred through the conduction or valance bands (non-localized part of the energy spectrum), as observed in the case of crystalline semiconductors. But hopping of electrons may occur from one localized state to another. The addition of conductive fillers in the polymer matrix culminated in the creation of localized states around the Fermi level which is located at the mid point between the bottom of the conduction band and the top of the valance band.6*36 According to the Mott-Davis 3 9 if the Fermi level lies in a band of localized states, the carriers (that is, electrons or holes) can move between the states by means of phonon-assisted tunnelling or hopping. With an increase in the filler concentration, aspect ratio, structure and particle size, the density of the localized states increases and the jumping distance between the two states decreases during hopping. Although hopping appears to be the predominant mechanism in carbon black-polymer composites, a transition from the hopping type of conduction to an effective band type conduction occurred when the density of the charge carriers was increased by applying a square law of DC voltage across the sample.40 The mode of conduction also changes from tunnelling to direct inter-aggregate conduction when the concentration of the carbon black in the polymer matrix increase^.^ At high concentrations of filler, direct inter-aggregate conduction is observed because of the presence’ of tighter aggregates in the polymer matrix. Thus, in conductive polymer composites at ambient temperature and normal mncentration of filler, two dominant mechanisms of DC conduction are observed, namely, electron hopping and a tunnelling mechanism. Although these two mechanisms may be experimentally indistinguishable, it was observed that hopping is the predominant conduction mechanism in the case of carbon black filled polymer composites, as observed by AC conductivity mea~urement.~’ The nature of the majority of the charge carriers, their concentration and drift mobility depend on the types of conducting fillers and their concentration. These

172

Short fibrepolymer composites

Table 7.1. Concentration (n) and drift mobility ( p ) of carriers in carbon fibre filled neoprene vulcanirates, at 25 - C

Conductivity Method of saniple preparation Composite no." Mill-mixed method Cement method

15m 20m 30m 5c 10c 15c

20c 30c

cr[(Q

m)- '3

0.72 6.5 28.6 6.9 108.7 185.2 217.8 434.8

Carrier Drift mobility concentration of carriers n(m p(m2V s 1

'

1.9 x 2.3 x 7.3 x 2.8 x 1.2 x 2.5 x 1.2 x 2.3 x

lo2' 1023

1024

10z4 1 0 2 ~

1025 loz6

2.37 x 1.77 x 2.45 x 10-5 1.55 x 5.66 x 4.63 x 1.45 x 1.18 x

a The no. indicates the concentration of carbon fibre in parts per hundred parts of rubber. (Source: Ref 26, reproduced with permission of the publisher, American Chemical Society, Inc.)

parameters can be studied by various methods, such as Hall effect, thermoelectrical power, diamagnetic susceptibility and magneto-resistivity coefficient measurements. In carbon black composites the majority of the charge carriers are holes (a hole is an electron vacancy carrying an equivalent positive charge) left by trapping of electron^.^' However, Burton reported that the dominant charge carriers are electrons.42 The galvanomagnetic characteristics of the carbon fibre-rubber composites can be determined by the Hall effect, which is widely used as one of the basic methods for determining the electron transport parameters of the material^.^^.^^ It was found that electrons (majority of charge carriers) are responsible for current conduction in these composites. With an increase in fibre loading, carrier concentration increases significantly, whereas the mobility of the carriers exhibits a very slow decrease (Table 7.1).26 The decrease in mobility at higher concentration and aspect ratio of carbon fibre is basically due to the creation of more scattering centres arising out of the increased rubber-fibre and fibre-rubber-fibre interfaces in the rubber matrix.25926 Electrical conduction in carbon fibre filled rubber composites is believed to take place by a hopping mechanism. The number density of localized states, through which hopping of electrons occurs, increases and the jumping distance between the two states decreases, with increase in fibre concentration and fibre aspect ratio. This is due to the presence of a large number of close-knit inter-fibre contacts in the rubber matrix, which would also result in an increased probability of electrons jumping from one state to another.

7.4 Effects of processing factors on electrical properties 7.4, I

Effect of dispersion

The processing factors significantly affect the electrical properties of polymer c o m p o ~ i t e sThe . ~ ~volume resistivity of the composites increases with increase in dispersion It was found that the milling conditions have a marked effect

Electrically conductive rubber and plastic composites

I73

on temperature dependence of volume resistivity of the composites. During remilling, the energetic agitation and thorough blending of the polymer-filler mix tend to break down the filler aggregates, resulting in higher resistivity of the composite. Binko and Vahala reported the effects of dispersion on electrical conductivity of the carbon black filled polyethylene composites near the threshold region.46 It is evident that when the time of mixing is double the material with the same concentration (critical) of carbon black near the threshold region, practically lost its conductivity. The effect of the method of addition of carbon black on electrical properties was studied. When carbon black was added to the powder polymer, the chain structure of blacks was degraded during the melting process because of the high viscosity of the polymer, which ultimately resulted in a non-conductive composite. However, when black was added to the molten polymer, the materials showed a readily detectable conductivity in the comparable concentration of the carbon black. Beyond the threshold region, the influence of addition of carbon black into the powdered or molten polymer on electrical resistivity was insignificant. Conductive composites made by solution mixing require a lower concentration of carbon black than the conventional mill mixing method.47 Agari et al recently reported the effects of dispersion on both electrical and thermal conductivities of graphite filled polyethylene composites. l 6 The composites were prepared by four methods, namely powder mix, solution mix, roll-milled mix and melt mix. The improved electrical conductivity of the composites follows the following order: powder mix > solution mix > roll-milled mix > melt mix. In the composite prepared by the powder mix method, the graphite particles were arranged to surround powdery polyethylene (honeycomb) structure and accordingly the formation of a conductive network can be easily achieved, which is not the case in other mixing methods. 7.4.2

Efect of mould pressure

The effect of mould pressure gives rise to anomalous behaviour on electrical resistivity. Increase in mould pressure increases the electrical resistivity due to the fracture of carbon black structure caused by greater expansion of rubber when the mould pressure is being released.48 However, it has also been reported that increase in pressure leads to the creation of a three dimensional 'cluster network' of conductive carbon black and the resistivity of the composites decrease^.^' Compactness of filler particles in the polymer matrix at high pressure decreases the volume rFiistivity of the composite signifi~antly.~,~' Nasr et a1 have reported that the effect of hydrostatic pressure during vulcanization tends to rearrange the carbon black particles inside the rubber matrix, which leads to an increase in cond~ctivity.~

'

7.4.3

Efect of vulconizotion

An increase in the degree of crosslinking reduces the volume resistivity of the

Short fibre-polymer composites 10’2

1 O’O 1o8 1o6

1o4 1o2 1 1 o-2 10-4

0

0.1

0.2

0.3

0.4

0.5

volume fraction of carbon fibre 7.2 Effect of vulcanization on volume resistivity: (a,) and (a,) composites prepared by thermovulcanization having LID N 25 (a,) and Y 100 (a,); (b,) and (b,) composites prepared by conventional vulcanization having LID-25 ( b , )and = 100(b,). (Sources: Ref 15 and 25 with permission of the publishers, Elsevier Applied Science Publishing Ltd and Society of Plastic Engineers.)

composites, owing to the formation of a three dimensional network structure in the vulcanizates. Using a coaxial electrode probe for measuring resistivity of the uncured carbon black filled rubber compounds, Boonstra” observed that the resistivity of the sample decreases with time, after one or two minutes of the sample being pressed. A review on electrical conductivity of polymers modified with conductive powders summarizes the effects of vulcanization on electrical resistivity of the composites of various polymer-filler system^.^' Increase in conductivity with vulcanization time in acetylene black-rubber composites is believed to be due to the formation of highly branched carbon structures in the rubber matrix. Studies on resistivity of the nickel filled epoxy resin during hardening (with diethylenetriamine) showed a fall in the resistivity in the first exothermic stage of hardening by three orders of magnitude, and in the second stage, an additional two orders of magnitude.47 Grinblat et a1 investigated the effects of vulcanization at different mould pressures and temperature^.^^ It was observed that both physico-mechanical and electrical properties of the rubber composites change significantly with the vulcanization system. In the case of carbon black filled styrenebutadiene rubber (SBR) composites, the changes in electrical resistivity during the formation of crosslinks have been studied by Nasr et al.” A major change in resistivity was observed after two minutes of vulcanization. Recently Jana and co-workers reported that the higher crosslinking density of the vulcanizates prepared by conventional vulcanization systems

Electrically conductive rubber and plastic composites

I75

register lower volume resistivity than those prepared by the thermovulcanization technique (Fig. 7.2). 5 . 2 5

'

7.5 Effects of polymer matrices on conductive network formation Surface tension of the polymer plays an important role in conductive network formation. Miyasaka et al showed that the volume fraction needed to form the conductive network tends to increase with an increase in the surface tension of the polymer. l 7 While studying the electrical conductivity of carbon black filled ethylene-vinyl acetate copolymers, it was observed that for the composites whose matricescontain more than 30 wt% of vinyl acetatecontent (VAc)no sharp break was present in the plot of conductivity versus carbon content and conductivity increased continuously with increasing VAc content.' Presence of polar groups in the VAc component caused good dispersion of the carbon particles in the VAc-rich matrices. In polymer blends the carbon black tends to concentrate in one phase or even move to the interface between the two polymers.54 Recently Amin et a1 have discussed the conduction mechanism of blended rubber composites on the basis of quantum mechanical tunnelling and hopping mechanisms and studied the electrical properties under different temperatures and pre-extension condition~.~

'

7.6 Effects of the types of filler, their geometry and morphology

7.6.I

Types of filler

Although particulate carbon blacks are the most widely used conducting fillers, graphite-based polymer composites have also been i n ~ e s t i g a t e d . ~Several ~.~? models have been proposed to study the variation of the resistivity of a composite and its carbon black content.' Radhakrishnan studied the effect of inert fiIlers on electrical properties of conducting polymer composites in order to improve the physical properties." It was observed that as the concentration of the inert filler was increased, the resistance of the composite increased, slowly at first up to a certain (inert) filler concentration beyond which there was a rapid increase in the resistance approaching z non-conducting state. Different types of met& in various geometrical forms, such as particulate, flake and fibre, have been used for making conducting polymer c o m p ~ s i t e s . ~ ~ - ~ l The formation of segregated networks in metal filled polymer composites depends on the concentration, size, shape and oxide content of the metal p a r t i ~ l e s . ~ ~A , review ~ ~ , ~of~ the - ~ electrical ~ conductivity of the metal filled plastics discusses the different theoretical models for predicting the critical metal c~ncentration.~ Lee et a1 have reported the electrical properties, such as resistivity and current-voltage characteristics of magnetite (ferrite powder) loaded polyethylene composites.68 Exfoliated graphite flakes of different aspect ratios have been used to make

I76

Short fibre-polymer composites

conductive polymer composite^.^^^^^ The advantages of Ltsing such polymer composites over metals include light weight, low cost and relative ease of processing. The use of short fibres of carbon, graphite and metal coated carbon in making Balta conductive polymer composites has been of recent interest.z5*z6.61*71 Calleja et ul reported that carbon fibres provide charge transport over a large distance and the conductive carbon black particles improve inter-fibre cont a c t ~ More . ~ ~ recently, the use of carbon fibre in making pressure-sensitive conductive nitrile rubber composite and EM1 shielding materials has been r e p ~ r t e d3-76 .~ 7.6.2

filler geometry and morphology

In particulate fillers the important geometrical and morphological aspects are the particle size, structure and porosity on which the formation of individual aggregates depends. It has been observed that smaller particle size leads to the formation of more conducting paths per unit v o l ~ m e . ~The - ~ individual ~-~~ aggregate size is proportional to the filler particle size. The smaller particles coalesce to form small aggregates, but their greater number results in a smaller gap between theaggregates, which increases theconductivity of thecomposites by formation of a conductive network. Another important effect of smaller particles is to retard the attainment of good separation dispersion. In particulate metal filled polymer composites the particle size ratio of the polymer to the filler and the shape of the filler particles affect the electrical resistivity s i g n i f i ~ a n t l y A. t~ ~ ~ ~ ~ ~ ~ higher particle size ratio (polymer to filler), the distribution is largely segregated, thereby forming metal-metal particle contacts at a very low concentration of the filler. The effect of filler structure on percolation threshold was investigated by Janzen." At a particular concentration of filler, a high structure carbon black filled composite shows higher conductivity than a low structure black filled composite and at the same particle size, porous and non-porous blacks lead to the formation of conductive and highly resistive composites, re~pectively.~.'~ The aspect ratio plays an important role in the case of flakes and fibrous fillers. In flake and fibre filled polymer composites the network formation depends mainly on the fibre aspect ratio and fibre distribution in the m a t r i ~ . ' ~ . ~ ~

7.7

Effect of structural deformation on electrical resistivity

The resistivity of electroconductive polymer composites is known to vary as a function of strain and time under static and dynamic condition^.^^ The effects of strain and rate of strain on resistivity depend on the concentration of the filler particle^.^^,^^ Under large tensile strains (above 30%) the resistivity usually increases at first as the strain is increased (due to the breakdown of carbon black structure), passes through a maximum, and then decreases (due to the alignment of the carbon black particles). Peter and Podoba have studied the effect of electrical conductivity of conductive rubber composites under conditions of

Electrically conductive rubber and plastic composites

I77

repeated compression and shear.86 It is observed that at the start of dynamic shear and compression tests, resistivity initially increases and then decreases, depending upon the magnitude of dynamic loading and the concentration of the filler. Recently, several researchers have investigated the effect of strain on electrical conductivity of the various types of carbon black loaded rubber composite^.^^*^^ The changes in conductivity have been explained as a result of the interaction between carbon black and rubber, and the effect of the occluded rubber. Under different strains, the effect of pre-extension on resistivity has also been studied. The changes in resistivity have been interpreted as due to either breakdown or re-agglomeration, depending upon the concentration of carbon black.

7.8 Hydrostatic pressure effect

on resistivity

Only a few studies have been made on the effect of pressure on resistivity. In the case of non-conductive and undoped rubber composites, negative pressure coefficients of resistivity were observed, showing an increase in electrical resistivity with increase in pressure.” However, conductive rubber composites with electrons or holes as the predominant carriers register positive pressure coefficients of electrical resistivity. Several studies have been reported on the pressure effect of the electrical resistivity of rubber composites with metallic and carbon black filler^.^^-^^ Two types of pressure-sensitive rubber composites could be obtained: one type has an on-off switch function, changing abruptly from an insulating state to a conducting state; while in the other type the resistance changes continuously as the applied pressure is increased. Metal particles are generally used in the on-off switch type and carbon blacks for the continuously variable type. Pressure effects on electrical conductivity in nitrile rubber composites have been reported recently where both particulate carbon black and carbon fibre are used as conductive fillers.73

7.9

Temperature effects on volume resistivity

The dependence of volume resistivity on temperature in conductive polymer composites had been studied by Aneli and T o p ~ h i s h v i l i .The ~ ~ temperature coefficient of resistance may be positive (PTC), negative (NTC) or zero, depending upon the concentration of fillers and their geometry, and the nature of the polymers and the fillers.’ The PTC of a composite is due to the uneven thermal expansion coefficient of filler and polymer, which leads to an increase in average inter-particle distance in the polymer matrix. In contrast, the NTC is due to the activation of the thermal emission of electrons from the filler particles, as observed in semiconducting materials. The temperature independence of volume resistivity has also been reported for highly filled polymer composites, where the direct inter-aggregated path for electrical conduction remains unaffected at high temperature^.'^*^^,^^,^^ The thermal mismatch between fillers and polymers at high temperatures 5 9 2 5 . 2 6 * 5

1979.94*95

I 78

Short fibre-polymer composites

causes an increase in inter-filler particle distance, thus making tunnelling or hopping of electrons more d i f f i ~ u l t . ~A t~low , ~ temperatures, ~.~~ conductivity of polymer composites loaded with carbon black seems to be governed by electron tunnelling or hopping between the carbon particles in the polymer matrix. However, at high temperatures the probability of tunnelling or hopping of electrons gets suppressed, owing to the increased inter-particle gap. On the other hand, an increase in temperature assists the other transport mechanism, known as thermal emission, which leads to an exponential rise in conductivity on heating. Depending upon the carbon black loading, the same polymer-filler system showed anomalous behaviour in temperature dependence of volume re~istivity.~~ Anomalous behaviour in temperature dependence of resistivity has also been observed in certain cases of polymer-filler systems, because of the phase change in the modified polymer.47 Studies on several thermoplastics filled with different types of carbon black showed a several-fold increase in electrical resistivity at temperatures near the glass transition temperature (T,) of the Further increase in temperature (above T,) caused a fall in resistivity due to the formation of an independent carbon black network as a consequence of enhanced motion of filler particles and lower viscosity of the polymer.'OO.'O'Recently, Bernhard has reported that an independent network of carbon black particles is formed due to the formation of 'flocculated' structures at elevated temperatures. lo'. The temperature coefficient of resistance in metal filled polymer composites may be positive (that is, metallic type conduction) or negative (semiconducting behaviour), depending upon the concentration and oxide level of the filler particle^.^ The heating and cooling curves in resistivity versus temperature plots showed hysteresis, which tended to disappear at high concentrations of metal particles, because of the efficient heat transfer through the metallic contacts. Gul et a1 have reported that the hysteresis loop appears because of the thermal mismatch between the metallic and polymeric phases.lo3 Increases in concentration and oxide level of metallic filler particles changed the temperature coefficient of resistance from metallic type (PTC) to a semiconducting one (NTC).'04.'05 The effect of temperature on volume resistivity of fibre filled rubber composites is a complex phenomenon, which depends on various factors such as degree of crosslinking of rubber, thermal expansion coefficient, thermal conductivity and temperature-dependent electrical performance of rubber and fibre, fibre aspect ratio and fibre c ~ n c e n t r a t i o n . ' ~It~ w ~ as ~ ~observed '~ that the variation of electrical resistivity with temperature shows positive temperature coefficient at lower fibre concentration and with increasing fibre concentration and fibre aspect ratio, the temperature coefficient of resistance shifted towards zero (Fig. 7.3 and 7.4).Is The increase in volume resistivity with temperature is due to the thermal mismatch between rubber and fibres which deters the easy hopping of electrons between the inter-fibre gaps. Hysteresis was observed during heating-cooling cycles. The amount of hysteresis was found to depend on the carbon fibre concentration, the fibre aspect ratio and the degree of crosslinking of rubber

I79

Electrically conductive rubber and plastic composites

E

E

1.2

1.o

r 0

0.8

0.6 0.4 0.2 0

20

40

60

80

temperature

100

120

140

('c)

7.3 Variation of volume resistivity with temperature for the composites prepared by mill-mixing method (m). (Source: Ref 15 with permission of the publisher, Elsevier Applied Science Publishing Ltd.) 20, 30, 40 are parts of carbon fibre per hundred parts of rubber.

in the composites. (Fig. 7.5).25The hysteresis gradually decreases with increased fibre concentration, fibre aspect ratio and degree of crosslinking.

7. I 0 Galvanomagnetic properties of conductive rubber composites The galvanomagnetic characteristics of the materials were determined by the Hall effect, which is widely used as one of the basic methods for determining the electron transport parameters of materials. Generally this effect can be used to determine the nature of the majority of charge carriers and their concentration and mobility. The principle of the Hall effect is that if a magnetic field is applied perpendicularly to the current direction passing through the sample, a transverse potential difference appears in the samples due to the spatial charge distribution, which is known as the Hall ~ o l t a g e . ~ "A' ~measurement ~ of Hall voltage for an n-type sample (where electrons are the majority of charge carriers) of known dimensions, current and magnetic field gives a value for the electron (carrier) concentration, n:

17.31

Short fibre-polymer composites

I80

-

0.45 I

E

E r

0.40

heating

- - - *- - -

0.35

-

0.30

.=

0.25

0

w.

.-> .-VY

0.20

2

0.15

4-

v)

Q)

E

-0

0.10

3

>

0.05 0

20

40

60

80

temperature

100

120

140

('Cl

Variation of volume resistivity with temperature for the composites prepared by cement-mixing method. (Source: Ref 15 with permission of the publisher, Elsevier Applied Science Publishing Ltd.)

7.4

20

40

60

80

temperature

100

120

140

('C)

Variation of volume resistivity with temperature for the composites prepared by cement-mixing method (TV indicates thermovulcanization and CV conventional vulcanization). (Source: Ref 25 with permission of the publisher, Society of Plastics Engineers.)

7.5

Electrically conductive rubber and plastic composites

181

where

n = electron concentration (m-3), I = current (A), B = magnetic field (Vsm-’) q = sign of charge, + or - (C = As), a = thickness (m), VH = Hall voltage (V). Since all the quantities on the right hand side of equation C7.31 can be measured, the Hall effect can be used to measure accurate values of carrier concentration (n). If a measurement of electrical conductivity is made concurrently with measurement of the Hall voltage, it is possible to determine the mobility ( p ) of the major carriers from the Hall coefficient and the volume resistivity. Since conductivity is given by

then

where (r

= conductivity

(W’ m - l),

p = volume resistivity (Qm), p = mobility (m’ V- s- ’) and R , = Hall coefficient.

Here, we have taken an n-type material (where electrons are the carriers). If we take a p-type material where holes (+e) are carriers similar results are obtained, but the sign of charge (q), the Hall electromotive force (emf) ( V H )and the Hall coefficient (RH) will be positive. Simultaneous measurements of the Hall effect and the volume resistivity were performed on the same sample to determine the electron transport parameters.26 Generally a fixed magnetic field of 0.51 ‘Twas applied for measuring the Hall effect. Since the values of magnetic induction B, current I and thickness are known durlng the vo!ume resistivity measurement, the value of carrier concentration n, and the nature of majority of carriers and their mobility p, can be determined by using equations C7.31 and C7.51. Hall effect studies at different temperatures in carbon fibre-rubber composites indicate that with increasing temperature the carrier concentration decreases along with decrease in conductivity of the composites, although the mobility of the carrier increases. It was also reported that the decrease in carrier concentration with temperature is much higher than the increase in mobility of the carrier

I82

Short fibrepolymer composites

Table 7.2. Temperatureeffect on concentration(n) and drift mobility ( p )of thecarriers in carbon fibre filled neoprene composite prepared by mill-mixed method (composite no. 20m)

Temperature T (T) 30 40 50 60 70 80 90 I00

Conductivity .[(Om)-’]

Carrier concentration n(m 3,

5.10 4.60 4.25 4.16 3.73 2.91 2.62 2.38

1.6 x 1 0 2 3 1.1 x 1 0 2 3 1.3 x 10” 0.98 x 10” 4.2 x 10” 3.2 x 10” 2.4 x 10” 2.0 x l o z 1

Drift mobility of carriers p(m2V-’s- I 1 1.99 x 2.6 x 2.04 x 2.65 x 0.55 x 0.58 x 0.68 x 0.74 x

10-4 10-3 10-3

lo-’ lo-’

(Source: Ref 26, reproduced with permission of the publisher, American Chemical Society, Inc.)

(Table 7.2).26 The slow increase in mobility of the carrier at increasing temperature supports electron transport via hopping between the localized states similar to that observed in various non-crystalline materials.39 However, the above effect becomes less pronounced with the increase of carbon fibre concentration and aspect ratio.

7.I I

EM1 shielding effectiveness

The phenomenal growth of the utilization of electrical and electronic devices in military, industrial, commercial and consumer applications has created a high level of electromagnetic pollution in the environment, causing a serious problem of EMI. EM1 consists of any unwanted spurious, conducted or radiated signals of electrical origin that can create a serious problem in the satisfactory operation of extremely sensitive electronic equipment, as found, for example, in surveillance, traffic control, missile guidance, marine radar, radar tracking and other similar instruments used in defence, space exploration, research, communication and medicine. To achieve the proper functioning of the electronic equipment, scientists have come forward with new types of shielding material, such as electrically conductive polymer composites, for electromagnetic compatibility (EMC) of the equipment. These composites can provide effective shielding against radiated interference. Electromagnetic or radar absorbing materials (RAM) have been studied for a long time for the purpose of radar cross-section reduction (RCSR) of warships, fighter aircraft, etc. Many absorbing materials are now being marketed whose basic ingredients include carbon fillers. One of the applications of such absorbing materials is in the construction of indoor microwave anechoic chambers. Normally, because of their bulky and fragile nature, these materials do not find much application in operational weapons platforms, but with the advent of improved processing techniques, carbon black fillers and other types of conducting filler-based elastomers and plastics composites can be used for such purposes.

Electrically conductive rubber and plastic composites

7. I I . I

I83

The nature of €MI

The effects of EM1 can range from minor nuisance to serious consequences. An example of a minor nuisance is interference (snow) appearing momentarily on a television screen. This is due to the operation of an electronic or electrical appliance such as a car ignition or a hairdryer. On the other hand, a serious consequence could result when an interfering signal disturbs the normal operation of medical electronic equipment being used to monitor the condition of a patient under intensive care. These examples are indicative of the nature and potential consequences of EM1 and the importance of effective control that is needed to achieve EMC. For an EM1 situation, three elements must be present: (a) a source(man-made or natural) of conducted or radiated electromagnetic waves (in general, every electronic device generates some degree of EM1 emissions); (b) a propagating medium through which the waves are transmitted; and (c) a receptor which suffers adverse effects from the received signals. lo’

7.I I .2

Definition of € M I shielding effectiveness

Electromagnetic radiation can be divided into near field and far field regions. In near field, the electromagnetic signal can be predominantly an electric vector or a magnetic vector, depending upon the nature of the source. In the far field, plane waves exist in which the electric and magnetic vectors have an equal ratio in phase and are orthogonal to each other. The plane wave radiation (far field) is of most concern in measuring shielding effectiveness (SE), defined as: SE = 1 0 l ~ ~PlTo PI

~7.61

where P, is the transmitted power, P,is the incident power and SE is expressed in decibels (dB). The plane wave shielding attenuation is the sum of three components (Fig. 7.6),74 namely, direct reflection (R),internal re-reflection ( B )and absorption ( A ) . SE can be expressed as the sum of the three components (in dB): SE=R+A+B

C7.71

7. I 1.3 €MI shielding effectiveness of conductive polymer composites

The main activity of an EM1 shield is to produce a barrier made of electrically conductive material that attenuates the radiated or conducted electromagnetic energy through reflections (direct, R and internal re-reflection, B) and absorption, A . Although metals can provide adequate shielding against EM1 problems, conductive polymer composites are being increasingly used owing to their advantages in respect to light weight, low cost, design flexibility, and versatile electrical and microwave properties.59,67*82996*108*109 The SE of the conductive

I 84

Short fibrepolymer composites

7.6

Mechanism of EM1 shielding effectiveness. (Source: Ref 74 with permission of the publisher, IEEE Publishing Service.)

polymer composites depends mainly upon the concentration and conductivity of the filler particles. In addition, it also depends upon the reflection and absorption coefficient of the given filler, and their distribution a n d geometry o r aspect ratio (LID),present i n the composite. Kortschot and Woodhams reported the effect of filler geometry o n SE a n d proposed a n empirical correlation between electrical conductivity and SE of the composites.’ l o A few studies on SE of the particulate carbon black filled polymer composites against electromagnetic radiation of different frequencies have been reported.”~82.9hE M C elastomers and plastic gaskets of low volume resistivity have been reported to be in use for EM1 shielding and hermetic sealing, wherein silver balls o r silvered brass balls are used as conductive fillers.’07.’ The EM1 shielding effectiveness, almost comparable to that of exfoliated graphite alone, has been observed in the frequency range of 0-1000 MHz, for the composites of exfoliated graphite filled polyester.” Ahmad et a / reported the poor frequency dependence SE of talc filled polypropylene composites.’ l 2 Metal fibre and aluminium coated glass fibre filled polymer composites have

’

I85

Electrically conductive rubber and plastic composites

70

1

I

I

I

I

1

Y

v) v)

W

c a, .-> Y

(GHz)

frequency

7.7 Shielding effectiveness as a function of frequency for the composites prepared by mill-(m) and cement-(c)mixing methods. (Source: Ref 74 with permission of the publisher, IEEE Publishing Service.)

been reported to show ineffective EM1 shielding materials in the X-band (8-12GHz) frequency range.' 1 3 , ' l4 Ahmad et al characterized the nickel coated carbon fibre filled polypropylene composite by measuring the insertion loss, return loss and the equivalent input impedance in the X-band frequency range.'15 The changes of the SE as a function of frequency are almost insignificant in this frequency range. The effects of thermal stresses and chemical exposure on the EM1 shielding of the conductive polymer composites of various polymer filler systems have been investigated.82q1'6q'l 7 It is evident that both the thermal stresses and chemical exposure significantly affected the bulk electrical conductivity as well as the EM1 shielding effectiveness of the composites. Inherently conductive polymers also function as effective shielding materials.' " Ruckenstein and Park reported the EM1 shielding of polypyrrole impregnated conductive polymer composites.' The inclusion of polypyrrole in a porous cross-linked polystyrene host polymer forms a conductive network and can be employed for electrostatic charge protection and EM1 shielding. However, the main disadvantages of such materials include lack of storage stability, high cost and poor processability. More recently, EM1 shielding effectiveness and return loss of the short carbon fibre-rubber composites have been reported at the different frequency ranges of

'

I86

Short fibrepolymer composites

3.5

I

I

I

I

1

3 .O h

-rn TI v) v)

0

-

2.5 2.0

t

, p

- - -d

c L

3

1.5

/:-w-

Q,

L

1 .o

-

r

0.5

8

5c 30m 1 oc 20c

I

30c

I

I

I

I

I

9

10

11

12

13

frequency

(GHz)

7.8 Return loss as a function of frequency for the composites, prepared by mill-(m) and cement-(c) mixing method. (Source: Ref 74 with permission of the publisher, IEEE Publishing Service.)

100-2000MHz and 8 - 1 2 G H ~ . ~It&was ~ ~ found that with increasing carbon fibre concentration in the composites, EM1 shielding effectiveness increased (Fig. 7.7)74 and the return loss (Fig. 7.8)74 decreased. Composites from the cementmixed method with a high fibre aspect ratio of 100 showed higher shielding effectiveness and lower return loss values than composites from the mill-mixed method, with low fibre aspect ratio of 25. The composite with high fibre concentration, high fibre aspect ratio and high composite thickness showed maximum shielding effectiveness and minimum return loss, which indicates that the loss due to the absorption increased. From the relationship between volume resistivity and shielding effectiveness it is evident that at a particular frequency, the shielding effectiveness depends not only upon the conductivity of the composite, but also upon the fibre aspect ratio and the composite thickness (Fig. 7.9).7437s A sharp decrease in volume resistivity was observed at a critical fibre 7.9 (opposire) Correlation between EM1 shielding effectiveness and electrical resistivity/CF (carbon fibre) concentration: (a) composites preparcd by mill-mixing method; (b) composites prepared by cement-mixing method. (Sources: Ref74 and 75 with permission of the publishers, IEEE Publishing Service and Chapman and Hall.)

50

I

I

I

n

m

Mill mixing

U Y

I

L/D-25

frequency: 2GHz - - - - - - - -

40

-

8GHz thickness: 3.5rnrn

.c Q,

0

1.7mrn

-

0

0

-

20 8

WJ

c

I-

-

v)

0

.

‘q

- - - - - - - - - - _- - - - - _ _ ‘a,---- - _ _ _-_- - - _ _- - - - _ _ -Q------

b

a

1o-2

-

1

Y

1o2

1

1o6

104

volume resistivity

(ohm m)

Cement mixing

L/D-100

frequency: 2GHz

--------

v) v)

8GHz

Q,

c Q,

1o8

--

thickness: 3.5rnm

0

1.7rnm

0

10-~

-

.

10 -

0 ’

-

I

I

I

I

I

1

1o2

104

1o6

volume resistivity

(ohm m)

I m v

1o8

I88

Short fibre-polymer composites

concentration, beyond which further increase in fibre concentration caused slow decrease in volume resistivity but a rather moderate increase in shielding effectiveness.

References I Blythe A R, Electricul Properties of Polymers, Cambridge University Press, London ( 1980).

2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35

Scarisbrick R M, J Phi’s D Appi Phys, 6 (1973) 2098. Bhattacharya S K and Chaklader AC, Polyrn PIrist Tcchnol Eny, I9 (1982) 21. Medalia A, Rubber Chem Technol,59 (1986) 432. Kirkpatrick S, Rev Mod Phys, 45 (1973) 574. Stauffcr D, Introduction t o Percolution Theory, Taylor and Francis, Philadelphia (1985). Fisch R and Harris A B, Phys Rev B, 18 (1978) 416. Bcllamy A, Robin P and Michcron F, C o n f I n t Cnoutch Prtris, 111-3 (1982). Clarke PS, Orton J W and Guest A J, Phys Reo B, 18 (1978) 1813. Balberg I and Bozowski S, Solid State Commun, 44 (1982) 551. Sherman S D, Middleman L M and Jacobs S M, Polym Eng Sci, 23 (1983) 36. Benguigui L, Yacubowicz J and Narkis M, J Polym Sci, Polym Phys Ed, 25 (1987) 127. Hsu W Y, Holtje W G and Barkley J R, J Muter Sci Lett, 7 (1988) 459. Carmona F, Prudhon P and Barreau F, Solid Sture Commun, 51 (1984) 255. Jana PB, Plast Rubber Compos Proc Applics, 20 (1993) 107. Agari Y, Ueda A and Nagai S, J Appl Polym Sci, 42 (1991) 1665. Miyasaka K, Watanabe K, Jojima E, Aida H, Sumita M and Ishikawa K, J M a k r Sci, 17 (1982) 1610. Sumita M, Asai S, Miyada N, Jojima E and Miyasaka K, Colloid Polym Sci, 264 (1986) 212. Bigg D M and Bradbury E J, ‘Conductive Polymers’, in Polymer Science and Technology, Vol. 15, Ed R B Seymour, Plenum, New York (1981). Janzen J, J Appl Phys, 46 (1975) 966. Kusy R P, J Appl Phys, 48 (1977) 5301. Janzen J, J Appl Phys, 51 (1980) 2279. Ezquerra T A, Martinez-Salazar J and Balta Calleja F J, J Muter Sci Lett, 5 (1986) 1065. Balberg I, Binenbaum N and Bozowski S, Solid State Commun, 47 (1983) 989. Jana PB, Chaudhury S, Pal A K and De S K, Polym Eng Sci, 32 (1992) 448. Jana P B, Chaudhury S, Pal A K and De S K, Rubber Chem Technol, 65 (1993) 7. Kusy R P and Turner DT, Soc Plast Eny J , 29 (1973) 56. Malliaris A and Turner D T , J Appl Phys, 42 (1971) 614. Yacubowicz J, Narkis M and Benguigui L, Polyrn Eny Sci, 30 (1990) 459. Polly M H and Boonstra B B S T, Rubber Chem Technol, 30 (1957) 170. Van Beek L K H and Van Pul B I C F, J Appl Polym Sci, 6 (1962) 651. Van Beek L K H and Van Pul’BI C F, Carbon, 2 (1964) 121. Sichel E K, Gittleman J I and Sheng P, Carbon Black-Polymer Composites, Ed E K Sichel, Marcel Dekker, New York (1982) Ch. 2. Pohl H A, J Polym Sci C, 17 (1967) 13. Mott N F, Adu Phys Mag Suppl, 16 (1967) 49.

Electrically conductive rubber and plastic composites

I89

36 Streetman B G , Solid State Electronics Devices, Prentice Hall, New Delhi (1983). 37 Mott N F, Philos Mug, 22 (1970) 7. 38 Overhof H, ‘Hopping conductivity in disorder solids’, in Advances in Solid State Physics, Vol. XVI, Ed J Treuch, Viewing Braunschwcig, Braunschweig (1976) 239. 39 Mott N F and Davis E A, Electronic Processes in Non-crystalline Materiuls, Clarendon Press, Oxford (1979). 40 Carley Read R E and Stow C D, J Phys D,2 (1969) 567. 41 Mrozowski S, Carbon, 9 (1971) 97. 42 Burton L C, ‘Electrical Resistivity Relaxation of Carbod black Filled Rubber’, Paper No. 80, presented at the meeting of the Rubber Division, ACS, Indianapolis, Indiana, May 11 (1984). 43 Norman R H, Conductive Rubbers and Plastics, Elsevier, London (1970). 44 Abdel-Bary E M, Amin M and Hassan H H , J Polym Sci Polym Chem Ed, 15 (1977) 197. 45 Agayants I M and Delektor AA, Int Polym Sci Technol, 14 (1987) T/18. 46 Binko J and Vahala M, Int Pulym Sci Technol, 14 (1987) T/71. 47 Slupkowski T, Int Polym Sci Technol, 13 (1986) T/80. 48 Boonstra B B and Dannenburg E M , Ind Eng Chern, 46 (1954) 218. 49 Thomson C M, Besuden T W and Beumel L L, Rubber Chem Technol, 28 (1988) 828. 50 Matsushita R, Senna M and Kuno H, J Muter Sci, 12 (1977) 509. 51 Nasr G M, Amin M, Osman H M and Badawy, M M, J Appl Polym Sci,37 (1989) 1327. 52 Boonstra B B, Rubber Chem Technol, 50 (1977), 194. 53 Grinblat M P, Vasil’eva E B and Rozova N L, Int Polym Sci Technol, 14 (1987) T/34. 54 S c a r A K, Rubber Chem Technol, 54 (1981) 820. 55 Amin M, Nasr G M and Sobhy M S, J Muter Sci, 26 (1991) 4615. 56 Buechs E, J Appl Phys, 44 (1973) 532. 57 Fox L P , R C A Rev, 39 (1978) 116. 58 Radhakrishnan S, J Muter Sci Lett, 4 (1985) 1445. 59 Bigg D M, Polym Eng Sci, 17 (1977) 842. 60 Bigg D M , Pofym Eng Sci, 19 (1979) 1188. 61 Martinsson J and White J L, Polym Compos, 7 (1986) 302. 62 Gurland J, Trans Metal1 Soc A I M E , 236 (1966) 642. 63 Mdkhopadhyay R, De S K and Basu S, J Appl Polym Sci, 20 (1976) 2575. 64 Bhattacharya S K, Basu S and De S K, Composites, 9 (1978) 177. 65 Nobile M R, Nicodemo L, Nicolais L, Egiziano L, Lupo G and Tucci V, Polym Compos, 9 (1988) 139. 66 Bridge B, Folkes M J and Jahankhani H, J Muter Sci, 25 (1990) 3061. 67 Li Lin and Chung D D L, Composites, 22 (1991) 21 1. 68 Lee SH, Heo G, Kim K H and Choi J S, J Appl Polyrn Sci, 34 (1987) 2537. 69 Wang Y S, O’Gurkis M A and Lindt J T, Polym Compos, 7 (1986) 349. 70 Foy J V and Lindt J T , Polym Compos, 8 (1987) 419. 71 Bigg D M Electrical properties of metal-filled polymer composites’, in Metul-jilled Polymers, Ed S K Bhattacharya, Marcel Dekker, New York (1986) 165. 72 Balta Calleja F J, Bayer R K and Ezquerra T A , J Muter Sci, 23 (1988) 1411. 73 Pramanik P K, Khastgir D, De S K and Saha T N, J Muter Sci, 25 (1990) 3848. 74 Jana P B, Mallick A K and De S K, IEEE Trans Electromag Compat, 34 (1992) 478. 75 Jana P B, Mallick A K and De S K, J Muter Sci, 28 (1993) 2097. 76 Jana P B, Mallick A K and De S K, Composites, 22 (1992) 451. 77 Verhelst W F, Wolthouis K G, Voet A, Ehrburger P and Donnet J B, Rubber Chem Technol, 50 (1977) 735.

I90

Short fibrepolymer composites

78 Sircar A K and Lamond T G , Rubber Chem Technol, 51 (1978) 126. 79 Abdel-Bary E M, Amin M and Hassan H H, J Polym Sci Polym Chem Ed, 17 (1979) 2163. 80 Ajayi J D and Hepburn C, Plast Rubber Proc Applics, 1 (1981) 317. 81 Das D, Basu S and Paul A, J Muter Sci, 15 (1980) 1719. 82 Bigg D M, Polym Compos, 7 (1986) 69. 83 Milewski J V and Davenport D E, ‘Short metal fibers and flakes’, in Handbook of Reinforcementsfor Plastics, Eds J V Milewski and H S Katz, Van Nostrand Reinhold, New York (1987) 103. 84 Kost J, Narkis M and Foux A, Polym Eng Sci, 23 (1983) 567. 85 Burton LC, Hwang K and Zhang T, Rubber Chem Technol, 62 (1989) 838. 86 Peter R and Podoba M K, Int Polym Sci Technol, 14 (1987) T/74. 87 Hashem A A , Ghani A A and Eatah AI, J Appl Polym Sci, 42 (1991) 1081. 88 Hassan H H, Khairy S A, El-Guiziri S B and Abdel-Moneim H M, J Appl Polym Sci, 42 (1991) 2879. 89 Aminabhavi T M, Cassidy P E and Thompson C M, Rubber Chern Technol,63 (1990) 451. 90 Nagata M, I n [ Polyrn Sci Technol, 13 (1986) T/44. 91 Kanamori K, Int Polym Sci Technol, 13 (1986) T/47. 92 Kakizawa K, Int Polym Sci Technol, 13 (1986) T/40. 93 Aneli D N and Topchishvili G M, Int Polym Sci Technol, 13 (1986) T/91. 94 Amin M, Hassan H H and Abdel-Bary E M, J Polym Sci Polym Chem Ed, 12 (1974) 2651. 95 Amin M, Nasr G M, Khairy S A and Ateia E, J Appl Polyrn Sci, 37 (1989) 1209. 96 Sakamoto R, Int Polym Sci Technol, 13 (1987) T/40. 97 El-Mansy M K and Hassan H H, Int Polym Sci Technol, 15 (1988) T/7. 98 Klason C and Kubat J, J Appl Polym Sci, 19 (1975) 831. 99 Ghofraniha M and Salovey R, Polym Eng Sci, 28 (1988) 58. 100 Voet A, Rubber Chem Technol, 54 (1980) 42. 101 Narkis M, Ram A and Flashner F, Polym Eng Sci, 21 (1981) 1049. 102 Bernhard W, Polym Eng Sci, 33 (1991) 1200. 103 Gul V E, Electrically Conducting Polymeric Materials, Khimya Publication, Moscow (1968). 104 Nicodemo L, Nicolais L, Romeo G and Scafora E, Polym Eng Sci, 18 (1978) 293. 105 Bhattacharya S K, Basu S and De S K, J Appl Polym Sci, 25 (1980) 111. 106 Van der Pauw L J, Philips Res Rep, 13 (1958) 1. 107 Violette J L N , White D R J and Violette M F, Electromagneric Compatibility Handbook, Van Nostrand Reinhold Company Inc, New York (1987). 108 White D R J, E M I / E M C Handbook Series 4, Dan White Consultants, Inc, Gainesville, Va (1971). 109 Bigg D M, Mirick W and Stutz D E, Polym Taring, 5 (1985) 169. 110 Kortschot M T and Woodhams R T, Polyrn Compos, 6 (1985) 296. 111 Duff W G, A Handbook Series on Electromagnetic Interference and Compatibility, Vol. 1, Interference Control Technology Inc, Gainesville, Virginia (1988). 112 Ahmad M S, Abdelazeez M K and Zihlif A M, J Muter Sci, 24 (1989) 1795. 113 Baker Z Q, Abdelazeez M K and Zihlif A M, J hfater Sci, 23 (1988) 2995. 114 Musameh S M, Abdelazeez M K, Ahmad M S, Zihlif A M, Martuscelli E, Ragosta G and Scafora E, Plast Rubber Proc Applics, 13 (1990) 237. 115 Ahmad M S, Abdelazeez M K, Zihlif A M, Martuscelli E, Ragosta G and Scafora E, J

Electrically conductive rubber and plastic composites

191

Mater Sci, 25 (1990) 3083. 116 Bigg D M , Polym Compos, 8 (1987) 1. 117 Osawa Z and Kobayashi K, J Mater Sci, 22 (1987) 4381. 118 Williams N, Varadan V K , Ghodgaonkar D and Varadan VV, I E E E Trans Electromagn Compat, 32 (1990) 236. 119 Ruckenstein E and Park J S, Polym Compos, 12 (1991) 289.

Design and applications of short fibre reinforced rubbers M C H LEE

8. I

Introduction

O n e of the most important areas in automotive engineering is to isolate the vibration and noise of vehicles effectively. Sources of vibration a n d noise include roadjtyre interaction, aerodynamic/vehicle interaction and power-trainjvehicle interaction. With the trend of making cars smaller, automotive isolation components o r systems become more important in achieving desired customer satisfaction. I t has been estimated that among the currently used isolation components o r systems, over 90% are classified a s passive components o r systems. Typical examples of passive isolation components o r systems for passenger cars a n d trucks are control a r m bushings, engine mounts, shock absorbers, steering linkage bushings, steering ball joint bushings, track bar bushings, strut rod mounts and torque rod bushings. The materials often used in fabricating passive components or systems are elastomer and thermoset composites, because these materials can be formulated and processed to have the required properties (e.g. high damping coefficient and high ultimate elongation) for effectively isolating automotive vibration and noise. Most passive isolation components use materials that have isotropic mechanical properties.' ' In other words, these properties d o not change with respect to the direction o f vibration and/or the component design. The isotropic performance characteristics may generate several application andjor performance deficiencies. For example, in a smaller vehicle, the spring rate of a n elastomer component needs to be reduced i n order to provide the required performance for isolating the vibration and noise of the car. However, the isolation components or systems made of a soft/isotropic elastomer material could lead to undesired yaw, roll and/or pitch motions of the car and consequently, could worsen the overall isolation capability, controllability and driveability. To overcome these concerns, automotive isolation components or systems have to be specially designed to provide the desirable anisotropic dynamic mechanical performance, However, these designs, such as metal insert with special configurations and protective metal shell, often increase the manufactur-

Design and applications of short fibre reinforced rubbers

I93

Table 8.1. Compositions of unidirectional short fibre reinforced polychloroprene composites Ingredients

Phr

Polychloroprene (Neoprene W) Magnesium oxide Zinc oxide Stearic acid Ethylene thiourea Dioctyl sebacate Sulphur Chopped glass fibre

100 4

5 0.5 0.7 12.5 0.5 0, 5, 10, 20

Table 8.2. Volume and weight concentrations of short fibre in polychloroprene composites Volume concentration (YO)

Weight concentration ( O h )

0 3.83 1.36 13.70

0 4.34 8.32 15.36

ing cost drastically. From the viewpoint of material science, elastomer composites filled with anisotropic fillers, such as short fibres, seem to be a feasible and cost-effective solution. Many researchers5-l have investigated the fracture mechanics of short fibre reinforced rubber composites. Little effort has been put into such research areas as the effects of fibre concentration and fibre orientation on the anisotropic modulus properties of short fibre reinforced elastomer composites. Therefore, in this work, we have investigated the modulus and tensile properties of a unidirectional short fibre reinforced polychloroprene composite system. The anisotropic mechanical properties of this elastomer composite system have also been determined.

8.2 Materials The elastomer used in this work was a polychloroprene(Neoprene W, Du Pont). Theshort fibre used was a chopped glass strand (PPG,Type 3075) with a nominal fibre diameter of 13.2 pm and a cutting length of 12.8 mm. The compositions of the short fibre reinforced polychloroprene composites investigated are shown in Table 8.1; the volume and weight concentrations of the fibres used in each of the elastomer composites are listed in Table 8.2.

8.3

Sample preparation

For each composition, mixing of the ingredients was carried out in a 1.2 kg Banbury internal mixer for 4min at mixing temperatures ranging from 30 to 110 "C. Each masterbatch was then mixed using a 152.4 x 304.8 mm two roll mill for 9 min.The short fibres in the matrix were successfully oriented unidirectionally (parallel to the rolling direction of the mill) by controlling the rotational

I94

Short fibre-polymer composites

force

force

e=oa

longitudinal

e=45*

e=go*

diagonal

transverse

8.1 The angles of fibre orientation in thc uniaxial tensile test.

speeds of the rollers (25 and 34rpm) and the gap distance between the rollers (1-3 mm). These unidirectionally oriented masterbatches were then moulded into standard slabs by compression moulding. The cure temperature and cure time used were 160"C and 25 min, respectively. Tensile specimens with three angles of fibre orientation namely, O", 45" and 90". were prepared by die cutting. A schematic description of the angles for fibre orientation in the tensile specimens is shown in Fig. 8.1.

8.4

Test procedure

Uniaxial tensile tests were conducted at room temperature using an Instron tensile tester. The crosshead speed chosen was 508 mm min- I which corresponds

Design and applications of short fibre reinforced rubbers

I95

to an initial rate of deformation of 0.33 s- l. Fracture morphology of each tensile specimen was determined using a scanning electron microscope (International Sci. Inst. ISI-DS-130).

8.5

Stress-strain behaviour of unidirectional short fibre reinforced elastomer composites

The stress-strain data of all the compositions at various angles of fibre orientation were analysed using a modified Mooney-Rivlin equation.l2 This equation has the following form: ln[o,/(h -

= ln[G(a))]

+ (l/h).ln[G(O)/G(co)]

C8.11

where oE = engineering stress, h = uniaxial stretch ratio, oJ(A - 1-') = reduced force, G(co) = shear modulus at (l/h) equal to zero, G(0) = shear modulus at (l/h) equal to one.

Equation [8.1] predicts that a plot of In[oJ(h - K 2 ) ]versus (l/A) will exhibit a linear relationship. This line can be extrapolated to intercept (l/A) at the initial shear modulus, G(O), and with (1/1) = 0 at the large strain shear modulus, G(oo), respectively. Equation [8.l] has been successfully applied to predict the isotropic stress-strain behaviour of various crosslinked rubber systems.1 2 * 1 3 In this section, we will show the stress-strain behaviour of the unidirectional short fibre reinforced polychloroprene composites investigated and their deviations from the stress-strain curves predicted using equation C8.11. The plotted results for the stress-strain behaviour of the polychloroprene composites with fibre concentrations of 13.70%, 7.36% and 3.83% (by volume) are shown in Fig. 8.2, 8.3 and 8.4, respectively. As shown in these figures, the stress-strain curves for the three polychloroprene composites with the angle of fibre orientation of 90" (with respect to the direction of uniaxial tensile deformation) exhibit linear relationships and can be described by equation C8.11. In each case, the data points deviate positively from the linear relationships (equation C8.11) at large values of stretch ratio, 1. This behaviour can be characterized by the mechanisms of strain-induced crystallization and/or nonGaussian network deformation of elastomers. 12-14 The stress-strain curves for the polychloroprene composites with the angles of fibre orientation of 45" and 0" clearly show the piecewise linear relationships at low and medium values of stretch ratio and also exhibit positive deviation from the linear relationships at high values of stretch ratio. The above piecewise linear relationships for the stress-strain data shown in Fig. 8.2-8.4 comprise two distinct regimes. In the first regime (at low values of stretch ratio), the value of o J [ ( h - h-')I is nearly constant with respect to the change in the uniaxial ratio. However, the value of aJ[(h - X -')I in the second regime (at medium values of

I96

Short fibrepolymer composites

m

P

Z

cc

0.4

A

N -

I

x I

o .2

x

Y

\

w

b

0.0

/ 0 0'

-0.2 0

0

45'

A

I

I

I

I

0.2

0.4

0.6

0.8

90' 1.o

l / s t r e t c h ratio 8.2 The stress-strain curves for the polychloroprenecompositeswith a fibre concentrationof 13.7% (by volume)and at various angles of fibre orientation.

stretch ratio) decreases with increasing stretch ratio and can also be described by equation [S. 1). The unique stress-strain behaviour for the unidirectional short fibre reinforced elastomer composites can be attributed to the oriented fibres in the elastomer matrix. The deformation mechanisms of this unique behaviour were determined with the aid of the fracture morphology of the specimens, and the results will be discussed in the next section.

8.6

Fracture morphology, fibre configuration and deformation mechanism

The fracture surfaces of the tensile specimens for all the fibre reinforced polychloroprene composites were examined by scanning electron microscopy. Figures 8.5, 8.6 and 8.7 show the fracture morphology of the polychloroprene composite with the fibre concentration of 13.7% at the angles of fibre orientation equal to 90", 45" and 0", respectively. The fracture morphology for each composite clearly shows an adhesional failure between the matrix and the fibres. This adhesional failure mode is typical for composite materials having a soft matrix and hard filler^.'^^^^^' The fracture morphology for the polychloroprene composite with 90" fibre orientation (Fig. 8.5) shows that most of the short fibres (over 80%) are unidirectional and parallel to the fracture surface. This result

Design and applications of short fibre reinforced rubbers

I97

0.6

0.4

0.2

0.0

-0.2

I 0

I

I

1

0.2

0.4

0.6

I

0.8

I 1 .o

1/stretch ratio 8.3 The stress-strain curves for the polychloroprene composites with a fibre concentrationof 7.36% (by volume) and at various angles of fibre orientation.

suggests that the deformation mechanism of the polychloroprene composites with 90" fibre orientation is the uniaxial stretch of elastomer among the parallel fibres. Consequently, the stress-strain behaviour of the polychloroprene compositions at 90" fibre orientation follows the modified Mooney-Rivlin equation for crosslinked rubbers (see equation [8.1] and Fig. 8.2-8.4 at the angle of fibre orientation equal to 90"). Above the value of critical stretch ratio of 3.0 (or the reciprocal value equal to 0.313) for the polychloroprenecomposites with different fibre concentrations, the stress-strain data for all the polychloroprene composites deviate pwitively '2-14 from the linear relationships described by equation

[S.l]. For the polychloroprene composite with the angle of fibre orientation equal to 0",the fracture morphology shows (Fig. 8.7) that most of the short fibres (over 80%) are perpendicular to the fracture surface. The deformation mechanisms for this case comprised two portions. The first deformation mechanism includes two steps, namely, the deformation of fibres (predominantly parallel to one another) and the deformation of elastomer matrix located among the ends of parallel fibres, and then followed by fibre pull-out/fibre fracture in the matrix. In this deformation mechanism the stress-strain behaviour shows a nearly constant value of oJ[(k - A-2)] at low values of stretch rstio (see Fig. 8.2-8.4). A critical

Short fibre-polymer composites

0.6

I

0.4 -

_ -

0.0

I

I

0.2

I

-0'

d

A

0

I

I

0

45O

A 90'

-

E-

0.4

0.6

0.8

1.o

l/stretch ratio 8.4 The stress-strain curves for the polychloroprene composites with the fibre concentration of 3.83% (by volume) and at various angles of fibre

orientation.

stretch ratio is therefore defined as the value of stretch ratio below which the first deformation mechanism is the controlling mechanism. Values for the critical stretch ratio of the polychloroprene composites with 0" fibre orientation range from 1.6 to 2.2 (or in terms of engineering strain from 60% to 120%) as the fibre volume concentration decreases from 13.7% to 3.83%. The values of the critical stretch ratio for the polychloroprene composites are much higher than those of the ultimate elongation for the glass fibres (the ultimate elongation being less than 0.050/,).This result is attributed to the deformation of the elastomer matrix among the ends of the parallel fibres before fibre fracture and/or fibre pull-out. The second mechanism is that after fibre fracture and/or fibre pull-out the uniaxial tensile deformation of the elastomer matrix dominates again, and therefore, the stress-strain behaviour of the polychloroprene composite follows that of the modified ivlooney-Rivlin equation (equation C8.11). A critical stretch ratio can also be defined as the value of stretch ratio below which the second deformation mechanism occurs. The values of the second critical stretch ratio for the polychloroprene composites with 0" fibre orientation and different fibre concentrations are about the same and equal to 4.0 (or the reciprocal stretch ratio equal to 0.25, see Fig. 8.2-8.4). Above this critical stretch ratio the stress-strain

Design and applications of short fibre reinforced rubbers

I99

0.5 mm 8.5 The uniaxial tensile fracture surface of the polychloroprene composite with the fibreconcentrationof 13.7% (byvo1ume)andat 90" fibreorientation.

curves deviate positively' '-14 from the linear relationship predicted by equation C8.11. The average length of fibres in the composite materials is always an important parameter which not only provides information on the reinforcement effectiveness of the fibres in a given matrix, but also provides information on improving the mixing methods for processing the composites. It is noteworthy that the average length of short fibres after mixing and moulding can be determined from the fracture morphology of the specimens with the angle of fibre orientation equal to 90" (Fig. 8.5). The average length of short fibres I,, determined from Fig. 8.5, is 268 pm; the average fibre diameter, d,, measured is 10pm. Based on these values, the average aspect ratio of short fibres lf/df, in the elastomer matrix was determined to be 26.8. As a reminder, the initial cut length of the fibres is 1.28 x 104pm. Using this value, in conjunction with the final fibre length in the matrix, the ratio of initial fibre length to final fibre length is 47.8. Thus, after Banbury mixing, two-roll mixing and compression moulding, the average fibre

200

Short fibre-polymer composites

0.5 mm 8.6 The uniaxial tensile fracture surface of the polychloroprene composite with the fibreconcentrationof 13.7% (byvo1ume)and at 45" fibre orientation.

length is reduced by 47.8 times from its initial fibre length. Even with such a severe reduction in fibre length, the aspect ratio of the short fibres in the polychloroprene matrix (lf/df = 26.8) is still much higher than the critical aspect ratio of the short fibres, lrc/df,required for having a maximum stress value15 l 7 (values of lfc/df being around 515). F o r the polychloroprene composite with 0" fibre orientation, such a high value of fibre aspect ratio also leads to the same deformation mechanism as described in the previous paragraph of this section for the first deformation mechanism of the polychloroprene composite with 0" fibre orientation. F o r the polychloroprene composite with the angle of fibre orientation equal to 45", the fracture morphology (Fig. 8.6) and the stress-strain behaviour (Fig. 8.2-8.4) also show a similar type of deformation mechanism as in the case for the polychloroprene composite with 90" fibre orientation. The value of aJ [(i- i.*)I for the composites with 45" fibre orientation is lower than that for the composites with 0'' fibre orientation. The deformation regime with a constant

Design and applications of short fibre reinforced rubbers

20 I

0.5 mm

8.7 The uniaxial tensile fracture surface of the polychloroprene composite with the fibre concentration of 13.7% (by volume)and at 0" fibre orientation.

value of o J [ ( i . - ;.-')I for the composites with 45" fibre orientation is longer than that for the composites with 0" fibre orientation. These findings can be explained as follows. In the case of 0" (45") fibre orientation the deformation of elastomer chains in the fibre-matrix interfacial regime can be simulated as the 180" (135") peeling deformation m e ~ h a n i s m . ' The ~ local stress (strain) of elastomer matrix in the 180" peeling deformation mechanism is higher (larger) than that in the 135" peeling deformation mechanism and, consequently, the fracture of fibres and the adhesional failure between the elastomer matrix and the fibres for the 180" peeling deformation mechanism are expected to occur earlier than those of the 135" peeling deformation mechanism. As a result, values for the ')I and the regime of stretch ratio with such a constant value constant o J [ ( i - iLof oE/[(iL - k-')] for the composites with 45" fibre orientation are greater than those for the composites with 0" fibre orientation.

Short fibrepolymer composites

202

8.7 Stress-strain equation for unidirectional short fibre reinforced elastomer composites

Based on the results discussed in the previous sections, a general schematic description for the anisotropic stress-strain behaviour of unidirectional short fibre reinforced polychloroprene composites was developed and is shown in Fig, 8.8. The corresponding stress-strain relationships at various angles of fibre orientation are expressed in the following mathematical forms. For 0" fibre orientation: (a) )iCl 2 h 2 1 ln[o,/(iL - >.-')I = ln[G(O,O")] (b) I,, 2 1, 2

[8.2a]

jLcl

ln[o&.

- l L - 2 ) ]= ln[G(co,W)]

+ (lLcI/i.).ln[G(O,O")/G(co,Oo)]

[8.2b]

For 45" fibre orientation: (a) I c 3 2 I 2 1 ln[oJ(i.

-

I-')] = ln[G(O,45")]

[8.3a]

(b) hc4 2 1, 2 h,, ln[aE/(h -

= In[G(co,45")]

+ (h,,/i.)~ln[G(O,45")/G(co,45")]

[8.3b]

For 90" fibre orientation: (a)

A,,

2i 2 1

ln[oJ(I - 1.K2)] = ln[G(co, 90")]

+ (i.c5/iL)4n[C(0,900)/G(co,90°)]

CS.41

The symbols in the above equations represent the same physical meanings as shown in equation [S.l]. The critical stretch ratios, i,,,i.,,, , Ic3, I,, and i.,,,are described in Fig. 8.8. These equations are useful for predicting the stress-strain values of unidirectional short fibre reinforced elastomer composites in automotive engineering design where numerical analyses, such as linear and non-linear FEA (finite element analyses) are required. Using equations [8.2a], [8.3a] and C8.41, we determined the values of initial shear modulus of unidirectional short fibre oriented polychloroprene composites at various angles of fibre orientation. In the next section, we will discuss the anisotropic modulus properties of the polychloroprene composites.

8.8 Modulus properties of unidirectional short fibre reinforced elastomer composites and applications

The values of initial shear modulus of unidirectional short fibre reinforced polychloroprene composites were determined from the stress-strain plots dis-

Design and applications of short fibre reinforced rubbers

0.8 I

-n

I

I

0

0'

0

I

45'

I

1

0.6

0.8

203

A 90'

0.6

CI

2

(u

0.4

I

x

CI

\

b" m

w

0.2

0

0 0

0.2

0.4

1 .o

1 /stretch ratio 8.8 The stress-strain curves of unidirectional short fibre reinforced elas-

tomer composites. cussed in the previous sections; the results are shown in Table 8.3. For all the three angles of fibre orientation investigated, the value of initial shear modulus increases as the fibre concentration increases. The effectivenessoffibre concentration on the reinforcement of the initial shear modulus of the polychloroprene composites is the highest (lowest) at the angle of fibre orientation equal to 0" (90"). Values for the corresponding relative initial shear modulus, G(0, qre1 (defined as the initial shear modulus of the composite with a given angle of fibre orientation, G(O,0),divided by the initial shear modulus of the elastomer gum, G(0)gu,,,)of the polychloroprene composites were also determined and are shown in Table 8.4. The plotted results for the initial shear modulus versus the fibre volume concentration are shown in Fig. 8.9. The results described in Fig. 8.9 clearly show that the modulus-composition equations suggested by Einstein' * and Smalland by Guth and Gold" cannot predict or describe the modulus properties of unidirectional short fibre oriented polychloroprene composites (except for the case of the polychloroprene composite with the fibre concentration of 3.83% and 90" fibre orientation). However, the relative modulus properties of the polychloroprene composites with different angles of fibre orientation can be successfully described by the thermodynamic relationship developed' for the modulus properties of multi-component polymer systems. This relationship is expressed in a general form as shown below:

Short fibre-polymer composites

204

Table 8.3. Initial shear modulus of uniderectional short fibre reinforced polychloroprene composites Initial shear modulus at various angles of fibre orientation (MPa) ~

0”

45”

90”

0.73

0.72 0.93 1.27 1.85

0.72 0.79 1.06 1.50

Fibre concentration ( ~ 0 1 % )

0 3.83 7.36 13.70

1.14

1.72 3.03

Table 8.4. Relative initial shear modulus of unidirectional short fibre reinforced polychloroprene composites Relative initial shear modulus at various angles of fibre orientation 0“

45“

90”

1.00

1.OO

1.56 2.36

1.29 1.76 2.57

1.00 1.10 I .47 2.08

Fibre concentration (~01%)

0 3.83 1.36 13.70

4.15

0

0’

4.0

-

3.5

- -------

3.0

-

2.5

-

A

45’

0

90’

0

Guth-Gold

1.5 2mo

i

0

Einstein-Smallwood

0

9 A n

2

4

6

Vf

8

10

12

[%I

8.9 Plots for the relative initial shear modulus versus the volumeconcentration of fibres of the polychloroprene composites at various angles of fibre orientation and predictions from Einstein-Smallwood and Guth-Gold equation.20

14

Design and applications of short fibre reinforced rubbers

205

5.5 v)

2

3

5.0

U

E

4*5

L

4.0

Q

= Q)

3.5

.-m

3.0

.r

2.5

v)

.-

CI

2.o 1.5 1 .o

0

2

4

6

8

10

12

14

8.10 Plots for the relative initial shear modulus versus the volume concentration of fibres of the polychloroprenecompositesat various angles of fibre orientation and predictions from the thermodynamic mixing equation.2

ln[G(O,8)"']

= +.exp[A(8)(4

- l)].lnK(8)

~8.51

In the above equation, A(8) represents the mixing index of the fibres with the angle of fibre orientation, 8, in the matrix; 4 is the volume concentration of the fibres; lnK(8) is the reinforcement effectiveness of the fibres with the orientation angle, 8; G(0,B)"' is the relative initial shear modulus of the composites with the angle of fibre orientation equal to 8. The value of the mixing index, A, of any type of filler in the matrix is non-negative.2 The higher the value of A is, the worse the degree of mixing of the fillers in the matrix.' The plotted results on the relative initial shear modulus as a function of the volume concentration of the fibres and the predictions frwn equation C8.51 are shown in Fig. 8.10. The values of 1nK and A determined for the polychloroprene composites with various angles of fibre orientation are shown in Table 8.5. Based on the results shown in Table 8.5, it can be concluded that values of reinforcement effectiveness, lnK, for the polychloroprene composites depend on the angle of fibre orientation, 8. The semi-logarithmic plot between the reinforcement effectiveness for the polychloroprene composites and the angles of fibre orientation exhibits a linear relationship (see Fig. 8.1 1).This linear relationship is expressed in the following mathematical form:

206

Short fibre-polymer composites

Table 8.5. Values of reinforcement effectiveness and mixing index for unidirectional short fibre reinforced polychloroprene composites Angle of fibre orientation, 0

Reinforcement effectiveness (InK)

Mixing index ( A )

14.58 9.41 7.12

0.33 0.36 0.33

0" 45"

90"

2.8

2.6

0

2.4

Y

c

2.2

2.0

1.8

'

0

I

I

1

30

60

90

fibre orientation

(degrees)

8.11 Semi-logarithmic plots of the reinforcement effectiveness versus the angle of fibre orientation.

ln[lnK] = 2.68 - 8.49 x 10-3.€l

(90' 2 9 2 Oo)

C8.61

The correlation ratio of the regressional analysis (Fig. 8.1 1) using equation [8.6] is 0.975. Contrary to the results of InK being a function of 9, the values of mixing index for the fibres in the matrix, A, are insensitive to the change in the angle of fibre orientation and are within the range of 0.33 to 0.36. This result means that the fibres are well dispersed in the polychloroprene matrix. As mentioned previously, the value of mixing index for a given filler in its perfect mixing state equals zero.2 In summary, for unidirectional short fibre reinforced polychloroprene composites, the value of 1nK decreases as the angle of fibre orientation, 0, increases. The value of the mixing index, A , on the other hand, is nearly independent of the angle of fibre orientation, 6. (This result is expected,

Design and applications of short fibre reinforced rubbers

207

Table 8.6. Initial shear modulus ratios of unidirectional short fibre reinforced polychloroprene composites Initial shear niodulus ratio Fibre concentration (vol%) ~~

0"/90"

45"/90" ~~

~~~~

0

~

.oo

1.01 1.44

1

3.83

7.36 13.70

I .62 2.02

I 20 1.23

1.18

since the mixing index, A, is a thermodynamic state variable that is only a function of such thermodynamic variables as surface Gibbs free energies for the fillers and matrix, volume fraction of the ingredients, temperature and pressure of the ~ y s t e m . ' ~As) a reminder, the values of both A and 1nK for elastomers filled with particulate filler (such as carbon blacks) are independent of the testing direction (or equivalently, the direction of deformation).2 Combining equation [8.5] with equation C8.61, we obtained the relationship for the initial shear modulus of the polychloroprene composites with various angles of fibre orientation. This relationship can be written in the following form:

+

ln[G(O,O)] = ~I[G(O)~~,,,] @exp[A.(+ - l)]~[lnK(Oo)]~[lnK(900)/lnK(Oo)~~900 [8.7]

In the above equation, all the symbols have the same physical meanings as discussed previously in the text. The range of the angle, 8, of fibre orientation is between 0" (longitudinal) and 90" (transverse) (see Fig. 8.1). The anisotropic characteristics of the initial shear modulus were described by twomodulus ratios as shown in Table 8.6. Each ofthe modulus ratios is defined as the value of 0" (45")modulus divided by that of 90" modulus. Values for both the 0"/90" modulus ratio and the 45"/90" modulus ratio increase with increasing fibre concentration. However, for the same fibre concentration, the value of the45"/90" modulus ratio is lower than that of the 0"/90" modulus ratio. The plotted results for the 0"/90" modulus ratio and the 45"/90" modulus ratio as a function of the volume concentration of the fibres are shown in Fig. 8.12. The above results are attributed to the different fibre orientations in the polychloroprene matrix. 8.9

Concluding remarks

In order to establish rules to assist the design of short fibre reinforced rubbers in various engineering applications, the mechanical properties and fracture morphology of unidirectional short fibre reinforced polychloroprene composites have been determined. Based on the findings, several deformation mechanisms for this composite system have been identified. The constitutive equations and several composition-processing-property relationships for predicting the stress-strain behaviour and the modulus properties of the elastomer composites have also been developed. The modulus properties of the polychloroprene composites with different angles of fibre orientation and various fibre concentra-

208

Short fibre-polymer composites

2.5

2.0

L

Q

1.5

1.o

0

2

4

6

8

10

12

14

8.12 The 0"/90" and the 45"/90" initial shear modulus ratios as a function of the volume concentration of fibres for the polychloroprene composites.

tions can be predicted by a thermodynamic relationship for the modulus properties of multi-component polymer systems. The above mechanisms and relationships are essential to the design and applications of short fibre reinforced rubbers. References 1 Lee M C H, J A p p l Polpi Sci, 29 (1984) 499. Lee M C H, Polyrti Eiig Sci, 25 (1985) 909. Trexler H E and Lee M C H, J Appl P o l p i Sci, 32 (1986) 3899. Lee M C H. J Appl Polyrtt Sci, 33 (1987) 2479. Thomas A G, Kiihher CIicw Techtiol, 48 (1975) 5. 6 Gent A N , Lindley P B and Thomas A G, J Appl PoIyrtt Sci, 8 (1964) 455. 7 Estes G M, Cooper S L a n d Tobolsky A V, J Mncroriiol Sci, Rev kJtrcromol Cherii, 4

2 3 4 5

(1970) 167. 8 Miwa M, Nakayama A, Ohsawa T and Hasegawa A, J .4ppl Polyrn Sci, 23 (1979) 2957. 9 Setua D K and De S K, J Mnier Sci, 20 (1985) 2953. 10 Akhtar S, De P P and De S K, J Mnrer Sci Leit, 5 (1986) 399. 11 Akhtar S, Bhowmick A K, De P P and De S K, J M n t r r Sci, 21 (1986) 41 79. 12 Lee M C H and Williams M C, J Polyrn Sci, Polyrii Phys E d , 23 (1985) 2243. 13 Lee M C H and Shen M, Polyrii J , 12 (1980) 495.

Design and applications of short fibre reinforced rubbers

209

14 Mark J E, Polym Eng Sci, 19(4) (1979) 254. 15 Lee M C H, unpublished research results. 16 Holister G S and Thomas C, Fibre Reinforced Materials, American Elsevier, New York (1966). 17 Kelly A and Tyson W R, ‘Fibre Strengthened Materials’, presented at the Second International Materials Symposium on High-strength Materials, University of California, Berkeley, California, June 17 (1964). 18 Einstein A, Ann Physik, 19 (1906) 289. 19 Smallwood H, J A p p l Phys, 15 (1944) 758. 20 Guth E and Gold 0, Phys Reu, 53 (1938) 322.

9 Design considerations and end-use applications of short fibre filled rubbers and thermoplastic elastomers A P FOLD1

9. I

Introduction

Nature has provided mankind with a vast array of composites, such as tree trunks, bamboo stakes and igneous rocks, which humans have for a long time put to good use. The realization that a combination of vastly different components can lead to a much improved or different kind of product led to the development of man-made composites, the best known early examples being rigid epoxy products reinforced with continuous strands or fabrics of glass fibres and bundles, mainly to provide impact resistance. Vulcanized rubber articles, reinforced with an array of particulate matters, such as carbon black or clay, were not recognized for quite a long time as being composites. Yet, for a material to be a composite all it needs is to consist of two or more distinct, separate phases with a distinguishable interphase between them, and some of its properties have to be radically different from those of its constituents. The main ingredient of a composite, the matrix, provides the basic characteristics; the other ingredients, which are present as a discontinuous elongated or particle phase, are called the reinforcers, if they provide property improvement, and fillers, if they do not. With short, individual fibres as the reinforcer, one can impart dramatic changes - even at low fibre concentrations - to the mechanical, viscoelastic, thermal and other properties of the compounded rubber. Excellent reviews have been published in the past few years on short fibre reinforcement of rubber.’-4 Some of this basic information, along with more recent findings, is presented in sections 9.2 and 9.3, because they are vital to the understanding of the proper employment of the reinforcer and offer guidance as to potential end uses. Recently, the focus has been shifted to thermoplastic elastomers as potential candidates for matrices. Unless otherwise cited, such thermoplastic elastomers are included in the term ‘rubber’ when describing general property trends in this chapter. When citing examples or results, the exact nature of the matrix is always given. The term matrix in this chapter is used to describe the fibreless compound, although this compound itself may be a composite of an elastomer matrix and other separate phases.

Design and end-use applications of short fibre filled rubbers

9.2

21 I

Background

Reinforcement of rigid plastic(thermosetting)materials with short fibres has been practised for quite some time. However, the adoption of this technology to rubbery (flexible thermosetting) and thermoplastic materials has been very gradual, mainly because of the success in using continuous cord reinforcement in most industrial elastomeric articles, imparting strength without sacrificing flexibility. While short fibre reinforced rubber combines the characteristics of the flexible rubber matrix and the stiffness and/or tenacity of the reinforcing fibre, short fibre reinforcement is insufficient to replace continuous cord reinforcement. Nevertheless, the improvements imparted by short fibres to elastomeric matrices are so substantial that short fibre reinforcement deserves detailed, in-depth discussion as a separate technology. Because short fibres can be incorporated directly into most rubber and thermoplastic compounds along with the other additives using standard rubber mixing equipment and because the resulting composites can be processed in standard rubber processing steps (extrusion, calendering, as well as compression, injection and transfer moulding), economical, high volume outputs are feasible. This is a significant advantage over the slower processes required for the incorporation and placing of continuous fibres, cords and fabrics. Thus, the penalty of sacrificing noticeable reinforcing strength with discontinuous fibres is counterbalanced by processing economies. In addition, short fibres significantly outperform simple particulate materials as reinforcers. The success of applying short fibre reinforced elastomeric materials lies in the understanding of the functions of their constituents, and the interaction between them. These will be discussed, rather succinctly, in the following sections before dealing with already proven and potentially useful applications. To utilize short fibre reinforcement efficiently, it is essential to have a working knowledge of (i) the properties of the matrix material; (ii) the relevant properties of the reinforcing fibre; (iii) the methods of uniformly incorporating the fibre into the matrix; and last but not least (iv) the intended performance(i.e. function) of the finished article. In many cases, short fibre reinforced articles fail prematurely because the designer did not consider the forces the fibres are exposed to during the operating life of the article. Because of this, a more than,general description is given with each new application about its function(s). The author assumes that the reader is fully familiar with the nature of the matrix to be reinforced. If not, abundant literature has been written on various rubber and thermoplastic materials. The author recommends two sources as a start: Goettler and Shen’s article’ and the Eliisromer Technology H i i n d h ~ o k . ~

9.2. I

Manufacturing constraints

Manufacturing constraints refer to those problems encountered during the process of incorporating the fibres into the matrix or which may manifest

212

Short fibre-polymer composites

themselves, to a different degree, during the final shaping of the end product. The primary concern is to prepare a fibre reinforced composite with improved properties. The improvements depend - besides the type of matrix and reinforcing fibre - on the uniform dispersion of properly oriented fibres with an adequate aspect ratio. Short fibre-elastomer composites can be prepared by most conventional rubber mixing compounding equipment: Banbury mixers, rubber mills, extruders, etc. The purpose of the mixing operation is to provide a homogeneous product in which not only the rubber ingredients but also the short fibres are uniformly dispersed. Good dispersion implies that the fibres are separated from each other (i.e. there are no clumps) and are surrounded by the matrix. Different mixing equipments do not produce composites with the same degree of uniformity and dispersion; Goettler and co-workers6*’found milling to be the best when comparing tensile results obtained with products prepared by various methods. Dispersing fibres in rubber is more difficult than dispersing the same amount of carbon black. The inherent nature of any non-rigid fibre tends to result in entanglement which is worse with increasing fibre length and increasing fibre concentration. Generally, it is recommended that fibre lengths be kept well below 25 mm and fibre loadings, especially with high performance fibres, below the 10 to 15 phr level. A few years ago, a fibre with highly fibrillated surface was introduced in the form of Kevlarb p-aramid pulp (see detailed item description in section 9.2.2). This fibrillated surface is of great benefit once dispersed in an elastomeric matrix but when attempting to mix p-aramid pulp directly into a rubber compound under moderate shear conditions, the fibrils preferentially cling to one another, creating clumps that are hard to disperse. There are several methods available to overcome, or at least mitigate, the fibre entanglement. They include pre-opening the fibres, pre-blending with other ingredients, using a concentrate, or adding the fibres at a viscous stage. Specialized mechanical devices (high speed mixers and pre-openers) separate the fibres in a dry form using high speed chopper blades. This method, by introducing a lot of air, decreases the bulk density. Owing to the greatly increased volume, this product cannot be shipped economically and should be used in situ. The earliest method used for disentangling the fibres was pre-coating them with a powdery and/or oily substance (usually one of the ingredients of the final matrix). This method does not require expensive or special equipment but it is also usually much less effective than the others in improving dispersion. A successful offshoot of this method is the preparation of fibre-thermoplastic mixtures where the thermoplastic is in a powdery form; usually mixing via tumbling is an adequate method to disperse the fibres uniformly. Preparing a concentrated elastomer-fibre masterbatch has proven very effective, not only from an economical point but also because the concentrate (which does not necessarily have to contain the same elastomer as the bulk of the matrix) disperses easily in most common elastomers and on all common rubber compounding equipment. This kind of masterbatch (generally in pellet or crumb form) is usually

Design and end-use applications of short fibre filled rubbers

213

'

c

CMD

common orientations

machine direction

orientations in abrasion testing

9.1 Definition of orientation directions.

marketed by the fibre manufacturer as a service to its customers. A variation of this method is the one used by some gasket manufacturers who disperse the fibre in a viscous elastomeric solution. If the already mentioned routes do not suffice, one should consider adding the fibre, gradually, at a viscous stage of compounding, because the higher the shear, the easier it is to obtain a uniform dispersion. More detailed description of these methods is given in Ref.4, p. 145-6. Orientation has crucial importance in fibre reinforcement as explained in section 9.3. Here, the definition, nomenclature and attainment of orientation will be discussed. There is a distinction between manufacturing and end use orientation; the former is introduced by the compounding equipment, the latter by the preparative technique used in making the end product. There are three orientation directions possible during manufacturing as shown on the left in Fig. 9.1. Machine direction or M D refers to the direction the rubber sheet is pulled from the calender or mill rolls. Since most of the fibres are going to be oriented in this direction, M D is also usually the same as fibre orientation. Cross-machine direction or CMD refers to the direction perpendicular to M D but in the plane of the rubber sheet. P M D is the direction perpendicular to both M D and CMD, and to the plane of the rubber sheet. (Some authors prefer to use other nomenclature: for instance, T for transverse = CMD, L for lateral = MD, and N for normal = PMD.) For certain end uses, namely those where abrasion resistance is the goal, there is a need for another, distinct orientation nomenclature. This direction is called on-end in our discussion and is shown on the right in Fig. 9.1. It basically implies a direction perpendicular to the plane of the rubber sheet. In normal rubber manufacturing processes, such fibre orientation cannot be achieved. Therefore, for abrasion testing, sample sections - cut out in MD - are rotated 90" and are moulded together to form the final test piece. Uchijayama et al devoted an entire paper' to friction and wear. They used an excellent graph to show the details of the test and the fibre orientation (their ' N direction corresponds to our 'on-end'). There has been sporadic confusion in the literature with orientation direction. In this chapter, the orientation designation always refers to the relationship of the

214

Short fibre-polymer composites

oriented fibre's direction to the direction of the applied force it is supposed to counteract. Unless noted otherwise, all observations, data, comments and recommendations pertain to composites which were tested and/or used with the fibres aligned in the machine direction, i.e. the fibres and forces were parallel. The efficiency (i.e. the magnitude of the effect versus the quantity of the fibre) of short fibre reinforcement depends on the aspect ratio of the fibre. This ratio, LID, compares the average fibre length to the average effective fibre diameter. Effective fibre diameter is the smallest width measurement of the cross-section. Commercial fibres have a very wide range of aspect ratios, from less than 50 to more than 3000. It is generally agreed that for good tensile reinforcement the aspect ratio should be at least 150. Aspect ratio can be easily changed by varying the fibre length. One could be tempted to use longer fibres to increase the aspect ratio. This, however, soon leads to processing difficulties. For optimum processing, staple fibre length should be between 5 and 20mm. Interestingly, transverse and shear moduli (both developing perpendicular to the fibre) were found' to be independent of fibre aspec. ratio. Brittle fibres (e.g. glass and asbestos) easily break up during high-shear p r o c e ~ s i n g , ' ~ ~thereby - ~ . ' ~ greatly reducing the residual, ultimate aspect ratio. Coran et all ' found that ribbon-like cross-section, instead of the usual round one, reduced fibre damage during processing because such fibres are more easily bent. Fibrillation is another damage which may occur during processing. It is typical of certain high strength. high modulus fibres (e.g. p-aramids) whose molecular structure resembles parallel rod-like assemblies. Fibrillation, however, results in an increase of the effective aspect ratio with its accompanying beneficial results. Actually, the highly fibrillating tendency of p-aramid is being used in making p-aramid pulp. Upon compression of the composite, micro-buckling can occur with stiff but ductile fibres, usually metallic ones, and it results in an irreversible orientation disruption, inasmuch as parts of the fibre will tend to respond differently to unidirectional strain. Temporary micro-buckling with non-rigid, organic fibres is easily removed when the fibres are subjected to a tensile strain exceeding the original compressive strain. 9.2.2

The fibre

There is a gamut of readily available fibres one could use for matrix reinforcement. These can be categorized according to their origin and physical characteristics. Chemically, they are inorganic and organic fibres, the former being of mineral (e.g. asbestos) or man-made origin (e.g. glass fibre). The organic fibres either occur naturally or are synthetically produced. The synthetic organic fibres were originally designed to imitate natural products (e.g. rayon) but in the past few decades the emphasis was placed on developing fibres with unique properties not found in nature (e.g. spandex). The advantage of man-made fibres over those found in nature is that the former can be produced with much narrower tolerances as far as properties are concerned. The other advantage is that

Design and end-use applications of short fibre filled rubbers

215

synthetic fibres can be produced in ‘infinite’lengths while nature’s fibres come in random and relatively short lengths (except for silk). This ‘infinite’ feature, however, is of no benefit when using short fibres. Therefore, many natural fibres can successfullycompete in non-demanding applications with the synthetic ones. As far as physical length is concerned, short fibres fall into one of the following categories, shown in order of descending length: staple, pulp, whiskers, fibrils and fibrids. There is no clear-cut demarcation line when one category ends and the other begins. Staples are usually greater than 5 mm long; fibrils and fibrids are microscopic. Staples are either precision cut, meaning that the vast majority is of equal length (with the remainingfew usually longer), or they are of random length distribution if they were produced via a stretch-break process. This latter version is usually cheaper to manufacture but it is also more difficult to use for short fibre reinforcement because it is more difficult to disperse evenly. As mentioned already in section 9.2.1, it is easier to disperse shorter fibres but the reinforcement efficiency is better with longer ones. Thus, one must strike a practical balance in selecting staple length. Pulp material is usually produced by some sort of grinding action of regular staple material. While the grinding itself causes randomization and reduction in the average staple length (e.g. 0.5-8 mm in Kevlara pulpI2),its main function is to provide a large array of curled, branched and often ribbon-shaped fibril-like appendages which are still attached at one end to the original fibre core. This arrangement provides mechanical locking to the matrix which can be most beneficial under compressive stresses, increases the fibre-matrix interface, and greatly increases the effective aspect ratio. Rigid, straight, inorganic fibres, usually with relatively low aspect ratio are called whiskers. Fibrils are basically diminutive staple with random length and cross-section distribution, and are usually the by-product of some mechanical abrasive force. Fibrids are similar but smaller than fibrils and are usually intentionally produced by precipitation under high shear. The actual choices of fibre are practically endless. Even within a single category, such as staple, one could find literature references - for practical or research end use - to nylon, polyester, polyvinyl alcohol, polyethylene (e.g. Spectra@),rn-aramid (e.g. Nomex”), p-aramid (e.g. Kevlaro), chemically treated non-regenerated wood cellulose (e.g. Santoweb”), asbestos, glass, wire, carbon, graphite, cotton, jute, silk, cellulose, coconut fibre, etc. Constant refinements are being made, for instance by increasing the aspect ratio (mainly by reducing the diameter) or by increasing the practical utility of the fibre. An example of the latter is Du Pont’s new KelvaP UltrathixTMpulp material’ which disperses more easily, mixes more smoothly and provides a more efficient thixotrope than standard p-aramid pulp. A revolutionary new idea was presented by Morgan14 that, by adding a special kind of fluoroplastic (not a fibre!) to the elastomeric mix, a discontinuous structure of short fibres and platelets was created within the elastomeric matrix during processing. The resulting composite showed improved tear resistance and abrasion resistance at both room and elevated temperatures. This kind of

216

Short fibrepolymer composites

reinforcement, owing to its novelty, is not included in further discussion on short fibre reinforcement. When choosinga fibre, one must be aware ofall properties associated with that fibre, not just those needed for the reinforcement of the matrix, but also those which may adversely interact with the matrix. An example, sometimes overlooked, is the moisture content of the fibre, either inherent or picked up during the manufacturing process, which may chemically interact with the matrix (such as with urethanes) or create voids or blisters in the finished product during thermoforming. Other aspects to consider are spin finishes and chemical residues which may have stayed on the fibre, or coatings or adhesives which have been intentionally introduced for a specific end use. 9.2.3 The matrix For the short fibre reinforcement technology, the following elastomeric matrices will be considered in this chapter: thermosets, thermoplastics, modified thermosets and thermoplastics. Thermosets vulcanized rubber - were the first matrices to be used with this technology, mainly in V-belts and tyre treads. A thermosetting elastomer develops a covalently bound, infusible and, usually, insoluble three-dimensional network after having been ‘cured’ from its original uncured state.” Curing of thermosets is typically carried out in the presence of catalyst(s) and/or heat. In the initial phase of the curing cycle, the material exhibits low viscosity and therefore is readily processable into a variety of shapes. The most common fabrication method of thermosets is pressure curing at elevated temperatures in a cavity mould. The extent of the chemical crosslinks determines the strength, rigidity, solvent resistance, as well as oxidative, chemical and thermal resistance of the elastomer. In the uncured, ‘green’ state, most thermosetting elastomers display a certain degree of tack (i.e. self-adherence). The term rubber is frequently used to describe a thermosetting elastomer. The most common rubbers are NR (natural rubber), IR (isoprene rubber), SBR (styrene-butadiene rubber), CR (neoprene = polychloroprene rubber) and EPDM (ethylene-propylene diene modified). Thermoplastic matrices were scouted because of the promising results with short fibre reinforced thermosets. Thermoplastic elastomers consist of high molecular weight linear or slightly branched polymeric chains which are solid at room temperature but can be fgrmed into a great variety of shapes at elevated temperatures and pressures. Thermoplastic elastomers exhibit the following positive features over thermosets: they offer potentially lower manufacturing costs (e.g. shorter processing cycles), they need less or no mixing with other additives, and they exhibit less batch-to-batch variability. Since thermoplastics can be remelted, remoulding, post-forming, joining and repairs can be effected via heating techniques.I6 On the negative side, they usually have no tack and require higher processing temperatures and pressures. The most commonly used ~

Design and end-use applications of short fibre filled rubbers

217

thermoplastic elastomers are polyester, polyamide and polyurethane-based, although recent studies17 indicate quite a bit of interest in polyethylene and polyvinyl chloride, most likely for economic reasons. The modifications of the existing thermosets and thermoplastics are carried out either to combine the beneficial characteristics of both (by blending or copolymerizing them together), or to impart certain, selected changes in specific properties. A typical example would be the use of XNBR (carboxylated nitrile) rubber as a reactive compatibilizer for blends of chlorinated natural rubber and epoxidized natural rubber, as reported by Ramesh and De." The use of these modified elastomers is quite new and not much has been scouted with short fibres. Whatever matrices are chosen for a given end use, they all fulfil numerous functions. They supply protection against abrasion and - often - against environmental influences (sometimes referred to as environmental corrosion), including ultraviolet light. They all provide stress transfer between the individual discontinuous reinforcing fibres and the applied forces and thus fairly well distribute the load in the composite. The matrix, especially in the event of highly oriented fibres, provides most of the 'transverse' (i.e. in the directions normal to the fibre orientation) durability both during extension and compression. In compression, the matrix supplies limited anti-buckling protection to fibres parallel or nearly parallel to the direction of compression. The extent of this protection depends also on the bending stiffness or rigidity of the fibres and on their aspect ratio. The short-term thermal resistance of a matrix can be increased greatly by the incorporation of strong, temperature resistant fibres. But the bulk of the composite's long-term temperature resistance depends on the thermal response of the matrix itself. On the other hand, fibres with poor thermal properties, especially with low melting point, will have a substantial, negative impact in the high temperature performance of the composite as the presence of fibres manifests itself as weak spots and/or voids. The choice of matrix, therefore, is very important. Goettler and Shen's' detailed summary of thermoplastic and heat-curable elastomers describes both common and specialty kinds. The reader is cautioned not to overlook the fact that the matrix properties depend not only on the elastomer component but on the combination of elastomeric and other possible matrix components: reinforcing fillers (e.g.carbon black), antioxidants, processing aids and oils, pigments, and - in the case of thermosets - accelerators, curing aids and crosslinking agents (e.g. sulphur). It is important to understand the chemical interaction between these components and the reinforcing fibres. Well-known examples are: the effect on certain p-aramid fibres of curing agents releasing free radicals, and the deleterious influence on PET (polyethylene terephthalate) of sulphenamide-type accelerators in the presence of moisture or of HRH (hydrated silica-resorcinol-hexamethylene tetramine) adhesion promoters. Also, as already hinted at in the previous section, surface treatments on the fibre or moisture therein can interact with the matrix ingredients.

218

Short fibre-polymer composites

9.3 TechnicaVengineering basics: effect of fibres o n matrix properties The information here is meant to illustrate the basic effects of short fibres in modifying the properties of the matrix, and the utilization of these changes for general kinds of end use problems. T h e examples in the following discussion use rubber as the matrix. Although the magnitude of the effect will be different with different matrix compounds, the trends are similar with thermoplastic elastomers. Here the discussion is organized according to properties and in the Applications section (9.4)the discussion is organized according to particular end uses. The two discussions will obviously overlap somewhat. A broader description of this basic information has been presented by the author in several media.4*’y , 2 0 After understanding the basics, the reader is urged to experiment with hislher matrix compound to select the best, the most cost effective fibre reinforcement. The various test methods for the characterization of fibre reinforced rubber were described by Abrate. Rubber World classified the tests normally used in the rubber industry in six basic categories2 stress-strain, time-dependent tests, destructive tests, durability, electrical properties and processability. Unless expressly stated otherwise, all reported dutu, coinmentS unil recornmentlutions in this chupter upply to composites which were tested untilor used with the fibres uligned in the muchine direction, i.e. the fibres were parallel to thefiwce they were supposed to resist. I t is ulso implied that the,fihres were evenly distrihuted in the matrix. Short fibres can reinforce both uncured (green) and cured rubber very effectively. In green rubber the following properties are of importance: yield strength which resists the tearing of the rubber during handling; lack of suficient modulus causes the rubber to sag on the mill rolls making handling difficult; too high cut resistance makes it strenuous to work (e.g. ‘cut a n d fold’) the rubber sheet on the mill roll and to remove the finished sheet at the end of milling; Mooney viscosity is a measure of how easily the green rubber will fill a cavity mould upon heating; tack, the ability of the matrix to stick to itself which is a n important asset when building u p an article from strips o r chunks of rubber; curing behuuioirr is usually described through viscosity changes - expressed by various torque versus cure time diagrams, e.g. Mooney scorch or oscillating disc rheometer. These diagrams predict how soon a compound starts to solidify after it has been subjected to heat and how soon it will reach its (nearly) fully cured state after which any further heating is counterproductive. Obviously, the last measurement has n o meaning for thermoplastic matrices which d o not undergo chemical curing. Most short fibres increase yield strength a n d the modulus very significantly. (Special methods had to be developed to measure stress-strain properties of green rubber having too much tack.”) The weaker the matrix, the greater is the relative reinforcement. The only type of fibre which did not exhibit improvement was short pieces of steel wire. Green cut resistance is increased to the point even with small loadings (e.g. 3 phr) that the mill rolls have to be stopped to allow the

’

’

Design and end-use applications of short fibre filled rubbers

.o

219

90

0 0

denim

4

+ 80 c

2

cotton floc

60 0

2

4

fibre content

8

6

10

(phrl

9.2 Erect of fibre type and content on Mooney viscosity.

worker to cut off the sheet laboriously, though safely. Mooney viscosity depends mainly on fibre content and somewhat on the nature of the fibre (see Fig. 9.2; data from the paper by Outzs2’). This is not the case with increased amounts ofcarbon black. If one were to use large quantities of carbon black to achieve equivalent modulus reinforcement in the cured rubber, the compound would have reduced tack, increased Mooney viscosity and much reduced scorch time. O n the other hand, short fibres barely influence the curing behaviour of the rubber. The effect of short fibres on tack has not been quantified, yet. The author’s personal experience is that tack is somewhat reduced in the presence of fibres above the 5-l0phr range. It is important to discuss the effect on cured rubber in more detail. The examples cited are with fibres which have no special surface treatment; no adhesive whatsoever was applied. The effects on stress-strain behaviour are mixed (see Fig. 9.3). Breuk elonyution is always sharply reduced but this is of no real concern in most applications where the article in normal use is not extended beyond 10%. Breuk strength is lowered i n the presence of fibres in most cases at ambient temperatures. Usually, the more a fibre reinforces the matrix (i.e. the steeper the initial portion of the curve = the higher the initial modulus), the lower the composite’s break strength. Figure 9.4 illustrates this point where the behaviour of six fibre types are compared with the fibreless control. Yet at elevated temperatures, break strength is usually higher with fibre, especially high melting fibre, in the matrix. Figure 9.5 describes the effect of increasing fibre loading with a given fibre. Note that with highly efficient reinforcers the stress-strain curve reaches a yield point quite early, after which a vacillating, up-and-down region appears. Mechanism-wise, at the yield point the phases start

Short fibre-polymer composites

220 20 1

I

I

I

I

I

I

1

control (no fibre)

10 v)

v) Q) L CI v)

5

0

50

100

150

200

elongation

250

300

350

1%)

9.3 Effect of p-aramid fibres on stress-strain behaviour. Fibre loading: 10 phr. Measured at room temperature in the machine direction.

to separate: the fibres are pulling away through the rubber matrix. It was established by microscopic examinations that even with weaker man-made fibres the fibres did not break in this region, nor did they at the point of ultimate composite failure. If one could economically and by practical means achieve chemical adhesion between the fibres and the surrounding matrix, the curve after the yield point ‘break’ would show a steady but slow increase until the ultimate break strength is reached. However, in many instances chemical adhesion is not necessary because in the applications cited later the fibres are oriented in such a manner that they - along with the surrounding matrix - have to extend. Upon extension, the surrounding matrix contracts perpendicular to the fibres’ axes, gripping them physically from the side. This phenomenon has been referred to in past literature as ‘the Chinese finger effect’. Until this point, the figures depicted behaviours when the samples were tested in the machine direction. It is very important to be always aware of the fibre’s orientation in the end use articles. Figure 9.6 demonstrates that 20% modulus (i.e. stress measured at 20% elongation), while showing a linearity for modulus versus fibre loading, is extremely influenced by fibre orientation (although not all the fibres can be oriented in a given direction with rubber manufacturing techniques and machinery). All other fibres show similar behaviour. Indeed, measuring a stress-strain property in the machine and cross-machine directions is the easiest

Design and end-use applications of short fibre filled rubbers

22 I

16

12

0

4

0 0

100 200 300 400 500

elongation (%) Effect of fibre type on stress-strain curve. 10 phr fibre in 50/50 NR/SBR: N6,6 =: nylon 6,6; E = polyester; G = glass; C = cotton; R = rayon; W = wire. 9.4

way to assess how effectiveone was in achieving directional orientation with most of the fibres.” Figure 9.7 shows the compressive behaviour of short fibre reinforced matrices. The diagram shows the test: as the cylindrical test sample is compressed to 60% of its original height, it becomes barrel shaped. Note that the fibres have to be oriented perpendicular to the compressive force to be highly effective. This also implies that the fibres in the mid-section of the ‘barrel’ are extended more than the ones close to the endplate and, therefore, carry the brunt of the load. There is a considerable increase in the force needed for compression with CMD, crossmachine orientation with increasing fibre loadings (as far as making the samples is concerned, the orientation is PMD since the cylinders are made up from thin discs cut out from the rubber sheet). However, if the fibres are aligned in the direction of force (these samples can be made up by extrusion or by rolling up a strip of rubber like a cigar), little additional force is needed beyond that needed to compress the matrix itself. High aspect ratio fibres, if compressed axially,

222

Short fibrepolymer composites

12

n

m

4

a

U

to v)

a

L c, v)

4

0 0

100

200

300

elongation 9.5

400

500

(%I

Effect of fibre loading on stress-strain behaviour of nylon 6,6 19mm staple in 70/30 chlorobutyl NR.

even if surrounded by a matrix, will either bend or permanently kink (e.g. wire) or shatter (such as glass). As they bend, they start to generate a vector component perpendicular to the compression force; hence the ‘little additional force’ mentioned above. With resilient polymeric fibre, the bending is not permanent: as the article regains its original shape, the fibres regain their original orientation. One would expect that, owing to the extensive elongation induced by the barrel-shaped distortion, there would occur an irreversible separation between the CMD oriented fibres and matrix. Hence, when the matrix retracts (upon removal of the compressive force) the fibres would be put under compression and, therefore, kink or crumple up. Subsequent compression then would be just enough to straighten out these fibres without extending them further: thus, no resistance would be put up by the fibres. This is, however, not the case. It was shown (see Fig. 6.11 in Foldi”) that repeated loading-unloading cycles do not cause any significant drop in the fibre’s reinforcing ability (measured at 60% compression and at a frequency of about 5 cycles/min).

Design and end-use applications of short fibre filled rubbers

0

5

fibre loading

10

223

15

(phr)

9.6 Effect of fibre loading and orientation on 20% modulus. p-aramid fibre in NK.

By now it should be obvious that the fibre’s ability to reinforce depends on its resistance to extension. This is also vividly demonstrated in measuring the bending stiflness of rubber bars reinforced lengthwise or crosswise with fibres. Figure 9.8 shows the diagram of the test piece, the applied force and the orientation of the fibres during bending. Considerable additional force has to be exerted to bend the bar if the fibres are lengthwise oriented on the extension side of the neutral plane of bending (in the illustration, that is the bottom of the bar). The fibres on the compression (top) side, since they are oriented parallel to the force being generated, add very little to the reinforcement. When the fibres are oriented side-to-side, i.e. parallel to the neutral plane and perpendicular to the arc of bending, there is no reinforcement whatsoever: the cords act more as a diluent. From this discussion, it is apparent that the greatest benefits are reaped from fibres located as far away as possible from the neutral plane of bending, and that the fibres on the compression side of the plane are just going along for the ride. Therefore, when producing an article which has to resist bending in one direction only, it is a waste of expensive material to put fibre-reinforced components on the compression side. However, if reinforcement is applied on the tensile side only, the neutral plane is moved from the centre toward the fibre-reinforced side, thus reducing the beneficial distance mentioned above. Nevertheless, putting fibres

Short fibre-polymer composites

base

0

I

I

I

I

10

20

30

40

fibre loading

50

(phr)

9.7 Effect of fibre loading on compression. Cylinder compressed to 60% of its original height; fibres oriented perpendicular to direction of compressive force; p-aramid fibre in NR.

where they do most good is a better practice than to put fibres into ‘sacrificial’ areas just to increase, marginally, the performance of the ‘beneficial’ fibres. In many industrial applications, the cover rubber (intended to protect valuable, continuous fibre reinforcement underneath) is subjected to cutting, gouging and similar abuses, making penetrution resistunce an important consideration. The author had to devise a test method to quantify the various components of penetration;” the test set-up is shown in Fig. 9.9. It was found that significantly more force was needed to move a four-sided pyramidal probe a given distance (19 mm) into the test sample if the fibres were oriented perpendicular (CMD) to the force of penetration (Fig. 9.10). However, considerable force (in addition to that needed to penetrate the unreinforced sample) was still needed to penetrate samples with fibres parallel to this force (MD). This seeming anomaly can be explained thus: even with the simplest of cutting shapes, there are three components of penetration. First, the object has to cut the surface. Since any object has a finite thickness, as it penetrates past the ‘wound’ it encounters two opposing forces: it has to overcome friction between its surface and the matrix

Design and end-use applications of short fibre filled rubbers

225

--~ 1

I

I

force

fibres extended on

fibres not extended

/

---------------D

0

E

0

5

10

fibre loading

15

20

25

(phr)

9.8 Effect of fibre loading and orientation on bending stiffness. Three-point bending; p-aramid fibre in 100phr NR; S = staple, P = pulp.

and it has to push (i.e. spread) apart the matrix material which is already behind its cutting edge. With the appropriate sequence of repeated penetrations, one can establish the contribution of each of these components (Fig. 9.1 1). As expected, more force (hence energy) is needed for wounding in the case of CMD oriented fibres. The extent of friction seems to be independent of fibre orientation (and, most likely, of fibre content). Spreading can be described as basically a series of bendings in case of M D orientation and a series of tearings with CMD orientation. Little has been said about tear resistance. The reason for this is that there are no unequivocal laboratory tests to predict the in-field tear performance of an article. Each laboratory test (and there are many) has its inherent flaws. They were originally designed to describe the tear performance of woven apparel fabrics (which they do well) and were adapted to coated fabrics first, then to mechanical rubber goods, and ultimately to fibre reinforced rubber. The ever-increasing proportion of elastomeric matrix to reinforcing fibre (from nil to 99) introduces a bias. For instance in ‘tongue tear’ (also known as ‘trousers test’), the more the reinforcing fibre (CMD oriented to the direction of the original tear pull) the lower is the tear resistance as measured by the energy needed for complete

226

Short fibre-polymer composites

force

I

sample

9.9 Schematic of penetration resistance test.

separation. The reason is that the tear, in the presence of large amounts of reinforcing fibres, does not propagate in the direction of the pull. Instead, it is channelled sideways, in the direction of the fibres. So, a tear which started as CMD may end up as a complete separation, but in MD. Finally, the subject ofabrasion has to be clarified. Using the National Bureau of Standards (NBS) abrasion test, it was found that the only fibre configuration that actually improved abrasion resistance was the ‘on-end’ orientation (Fig. 9.12). The other two orientations, i.e. when the fibres were oriented in the direction of rubbing or across it, were harmful to abrasion induced performance. The theories introduced by Thomas and Southern23 and S ~ h a l l a m a c hseem ~ ~ to be in agreement with these findings.

9.4

Applications: proven and potential

It is assumed that the reader is familiar with the basics; nevertheless, references are given for those needing a ‘refresher’ in the particular technology or application. The scope of this chapter does not permit giving blueprints and/or exact formulations for every application listed. In each case, however, the known critical parameters for the particular application are identified to allow the reader to predict how a given system may perform for an intended end use. The majority of the applications discussed fall into these broad categories: belts, hoses, friction products and automotive applications (mainly in tyres), representing a gamut of

227

Design and end-use applications of short fibre filled rubbers

1600

I

I

- - - - - - force

z 1400 -

I

I

perpendicular t o fibres

---

- -_-. . , ,,

Y

Q,

0

1200

-

,

P

_ - - -_ _ _ _ - - - - - - - -

CI

-

0

0

I I

c

I

1000

0

-

,/ ,,

I8 #

c.

,

0

I 8

0

-

s

Q)

c Q,

n

800

600

0

10

20

fibre loading

30

40

50

(p.hr)

9.10 Effect of fibre loading and fibre orientation on penetration:four-sided, pyramidal probe with 15" apex; 19mm penetration; p-aramid fibre in NR; S = staple, P = pulp.

general end use functions, such as power transmission, movement of material and reduction of viscosity or elasticity.

9.4. I

Power transmission belts

Power can be transferred either via solid (rigid) components or by the use of elastic ones. The advantages of the elastic ones are less noise, less mass and greater tolerance for shock load. The main disadvantage is that most elastic assemblies exhibit slippage, with timing be!ts being the exception. Gears, clutches or drives are used to transmit power between rotating shafts. Clutches transmit rotation at equal rotational speeds; gears and drives are usually designed to operate at different rotational speeds. Drives are either chain drives or belt drives. Because of their nature (e.g. construction material), chain drives and standard gear drives are outside the scope of this discussion; a specific use for a clutch assembly is discussed in section 9.4.4. There are three kinds of power transmission belts where short fibre reinforcement can improve performance characteristics: V-belts, flat belts and timing belts (see Fig. 9.13 for examples), of which V-belts are by far the most important. In

228

Short fibre-polymer composites

B

I probe

I sample

foundin!

friction

ioundin friction

preadin!

preadin

CMD

MD

fibre orientation 9.11

Effect of fibre orientation on the components of penetration.

each instance, the fibre’s ability to reduce bending, compression, wear and/or elongation of the matrix is utilized for the reinforcement. V-belts and flat belts rely on friction between the pulley (the term pulley here includes flangeless cylindrical drum) to convey power from one pulley to another. Idler pulleys may be included in the system but their only purpose is to change the direction of the motion, to support the weight of the belt or to maintain a steady belt tension. With a V-belt - which actually has the shape of a truncated V or trapezoid - the force is transmitted on its two non-parallel sides, while with flat belts - with an elongated rectangular cross section -friction is generated only on one side. The advantage of the V-belt design is that the normal force pressing against the sidewalls is much greater than the force pushing the belt into the pulley’s groove due to the geometry (see Fig. 9.14 for the vector analysis). With flat belts no such vector enhancement is possible. To transmit more power, one has three choices: increase the contact area, increase the tension or increase the coefficient of friction. With flat belts, increasing the contact area is achieved by increasing the width. Spatial limita-

Design and end-use applications of short fibre filled rubbers

140

I

I

I

I

229

I

E

-

120

100

cross machine

-

80 machine direction

MD CMD

60

40

I

I

1

fibre loading 9.12

I

-

I

(phr)

Effect of fibre loading and orientation on abrasion NBS abrasion test; p-aramid pulp in EPDM/neoprene rubber.

tions imposed on flat belts led to the development of V-belts where a more compact design offers more efficient power transmission. Increased belt tension necessitates the use of stronger reinforcing cords/cables, such as steel, glass or aramid. The coefficient of friction depends on the surface characteristics of the belt and can be changed with proper choice of matrix compound (for the basics, see Ref 5). 9.4. I . I V-belts As mentioned earlier, the angled sidewalls of the V-belt result in much greater

normal forces on the sidewalls of the pulley than the actual tension forcing the belt into the pulley. The smaller the belt angle, the greater the enhancement factor. For instance, at 60 O the normal force is twice that of the belt tension; at 26 O the normal force is 4.5 times. However, increased pressure also results in increased wear, in overloading the cord layer, and, with narrower grooves, in less tolerance to misalignment. A compromise of 40 O was reached internationally for a standard V-belt (and pulley groove) angle. It is necessary to point out that with a true V-belt the angle formed by the two sides is somewhat greater than the angle of the pulley groove. Therefore, when a V-belt is forced - wedged - into the groove, especially at high tension and during shock loads, a distortion of the originally flat reinforcing cord layer takes place: the cord layer becomes

230

Short fibre-polymer composites

FLAT BELTS

standard f l a t belt

multiple-V f l a t belt

V-BELTS

raw- edge standard V-belt

wrapped narrow V-belt

notched V-belt

multiple V-belt

SYNCHRONOUS or TIMING BELT

9.13

Power transmission belt types.

Design and end-use applications of short fibre filled rubbers

23 I

c

a,

n

9.14

Force enhancement in a V-belt: vector geometry.

V-belt

in use

neutral plane

properly loaded belt 9.15

overloaded belt

Distortion and load shedding in overloaded V-belt.

concave (Fig. 9.15). The result is that the cords on the edge of the belt are following a longer path than the ones in the centre, and therefore carry more of the load (i.e. the cords in the centre will ‘shed’ some of the load). In the past, different rubber compounds were used in the V-belt body to reduce this problem of concaving. A typical V-belt is constructed from many parts (see Fig. 9.16 for details). Sometimes, the entire belt is wrapped with an elastomerimpregnated fabric (such as shown in Fig. 9.13) to reduce the problem of cord ‘pop-out’ at the belt’s edges and, presumably, to increase both friction and belt life at the same time. These belts, owing to their increased labour demand (they have to be pre-shaped and wrapped individually, then cured in grooved, segmented

232

Short fibre-polymer composites

neutral plane

-outer c o v e r fabric rubber r e i n f o r c e m e n t c o r d reinforcement rubber cushion s t o c k base rubber inner c o v e r fabric

9.16 Typical V-belt constrtiction

moulds), are rapidly disappearing from large volume applications. Instead, raw-edge belts (i.e. belts which have been cut into the V-shape from a wide, cured sleeve) are now forming the bulk of the market. The t o p rubber layer and the optional t o p cover fabric are subjected to extension every time the belt passes over the sheave. On the other hand, the components below the cord layer that form the ‘ncutral plane’ go into periodic compression. The rubber along the sides is exposed to periodic shear. All this distortion gives rise t o considerable heat build-up. Thus, any efficient reinforcement has to be able to withstand not only the mechanical but also the thermal demands. Fibre reinforcement in the cushion stock immediately below the cord layer has to be oriented perpendicular to the cords (and parallel to the neutral plane) to be effective: this way, both the increased tenacity and bending stiffness of the composite are utilized. Such reinforcement prevents individual cords from cutting into the cushion stock, increases side-to-side bending stiffness (thus reducing the tendency for concavity) and also increases wear resistance along the contact areas with the sheave. Contrary to rubber compounds made stiff with high carbon black loadings which add considerably to heat build-up, fibres so oriented add very little to heat generation. The fibre loadings (between 3 and 5 phr) needed for significant performance improvement are both economical and practical when compared with the high carbon loadings (60-100 phr) required to achieve on 1y m it rg i n a 1 res LI 1t s. There is it great array of choices even for simple V-belts: classical, narrow, notched. Special, wider designs are used for variable speed belts. In-depth discussion of all sort of belt designs and functions can be found in the literature.15 Goettler and Swiderski,28give a good description of using cotton, polyester and short cellulosic fibres in a neoprene composition. With V-belts, the use of multiple, independent belts was the first step to higher power transmission capabilities. O n e drawback of this approach is that if all the belts are not exactly of the same physical dimension and/or of the same

’’

Design and end-use applications of short fibre filled rubbers

9.17

233

Photograph of a CVT belt.

construction, then some belts are going to carry more load than the others. This means either that the belt array has to be overdesigned, or that one of its members is going to fail prematurely. A recent development is the multiple V-belt (also called banded or joined V-belt) where several V-belts of identical construction are linked together at the top (like this: M ) by the matrix and a common fabric layer (separate from the reiaforcing cord layer). Since not much side force is generated in this common fabric layer, fibre reinforced sheeting could replace it. The fibre orientation should be across the circumference of the belt to utilize the increased modulus and tensile strength. It is imperative that such belts not only do not bottom out in any one of the individual sheave grooves but also the sheave design must be such that the belts do not touch the outer tip of the grooves. Otherwise, the wedging action - the salient feature of the V-belt concept - is lost and the belt becomes a very inefficient multiple-V flat belt. A very recent spin-off of V-belt technology is the development of special CVT (continuously variable transmission) belts which operate with ‘push’ rather than the ‘pull’ used by all other power transmission belts. While the general belt cross-section is still a truncated V (to take advantage of the wedge action), the belt is a series of hard rubber blocks individually attached to both sides of a tensile band. Some configurations have steel inserts to help in keeping the rubber blocks in place (Fig. 9.17). Reinforcement of these rubber blocks with fibres, both to increase hardness and to reduce wear, could be advantageous. This technology is too complex to go into detail here. Numerous automobile and belt companies (e.g. Fiat, Volvo, GM, Ford, Borg-Warner, Gates, Dayco) have developed and experimented with such belts. Most early publications were news releases from these companies. For general information on the use of fibres in V-belt compounds see the article by Rogers.*’

234

Short fibre-polymer composites

9.4.1.2 Flat belts Application of short fibres in flat belts is not as advantageous as in V-belts. The reinforcing cord layer does not get distorted on the drum, since there is near perfect mating between the belt and the drum. The only place where short fibre reinforcement of the matrix should be considered is the layer immediately below the cord layer. Using fibres oriented in the direction of travel would add only marginally to the overall modulus of the composite but it would contribute significantly to the bending resistance around small diameter drums. In the cross direction (i.e. perpendicular to the cords) there would be a modulus increase in a direction where it is not needed. While the bending resistance would not increase, wear resistance could be impaired. If the belt-reinforcing cord seems to bite into the underlying elastomeric matrix, either the belt load is too high, or the matrix should be recompounded for greater hardness. Only as a last resort is the use of fibre reinforcement recommended. The only really beneficial fibre direction is the one perpendicular to the belt’s friction surface. While this results in improved wear, it can be achieved only with small belts where the labour-intensive application of individually laid down strips of rubber could be justified. As mentioned earlier, one of the choices to transmit more power is to increase the contact area. One way to increase the contact area with a flat belt, without increasing its width or increasing the diameter of the drum, is to shape the side in contact with the drum like this: /”. In such multiple-VJat belts (also called Poly-ribTMbelts) the angle of the Vs and the V-shaped grooves in the drum are identical and the belt completely bottoms out in the grooves. Therefore, there is no wedging effect such as in the case of the multiple V-belt. The important differences between a multiple V-belt and a multiple-V flat belt are: the multiple V-belt has truncated rib-like components and the bottom of the belt does not touch the (usually rounded) bottom of the groove. A multiple-V flat belt has fully V-shaped ribs which completely bottom out in correspondingly shaped grooves with the reinforcing cord layer running uninterrupted along the entire width of the belt, while in some banded V-belts each individual truncated rib has its own, individual, cord layer reinf~rcement.~’ 9.4.1.3

Timing belts

Timing belts (also known as synchronous belts or synchro-cog belts) transmit power through a set of rectangular, trapezoidal or semicircular teeth which mesh with correspondingly shaped grooves on drums, usually with flanges on both sides to prevent the belt from ‘wandering’ off the drum’s surface. The teeth or cogs are joined to a reinforcing layer (usually fabric or unidirectional strands) of high modulus. There are two kinds of use for timing belts. One requires great precision (and therefore is also called precision drive belt) but only limited force transmission and it is found mainly in instrumentation and office equipment. The other is used as a machinery component, in lieu of chain drives, where frictional losses (unavoidable with V and flat belts) cannot be tolerated. A typical example is timing belts in today’s automobile engines.

Design and end-use applications of short fibre filled rubbers

235

The first kind of timing belt usually transmits marginal power loads and operates in very restricted spaces, the main function being the maintenance of synchronized motion between two rotating objects. The bulk of the elastomeric material is located in the teeth. Therefore, short fibre reinforcement in the belt tension member is not necessary nor is it feasible because of the very small amount of matrix surrounding the directional reinforcer. Usually the entire tooth surface is covered with a thin, tightly woven fabric, embedded in the matrix. With such belts fibre reinforcement could be indicated in the unlikely event of tooth loss during normal operation. The reinforcing fibres should be located as close to the teeth’s surface as possible, oriented in the direction of movement. A thin fibre reinforced layer of the matrix material could be placed in the mould between the outer fabric and elastomer forming the teeth. However, because of the extensive flow during tooth formation, one should expect certain orientation distortion in the finished product. The purpose of the fibres, oriented in the hoop direction, is to provide additional modulus increase at the stress areas where the cogs meet the reinforcing layer. In timing belts, which are meant for transmission of greater force, tooth breaking or gouging is more frequent. These belts are thick enough between the high modulus continuous reinforcing layer and the teeth to allowjudicious use of fibre reinforced laminates there. However, similar fibre disorientation should be expected during the moulding operation as mentioned above. The purpose of the fibreis similar to that mentioned for the precision belts: to reinforce the stress area where the cogs meet the continuous reinforcement. Note that sometimes it is more advantageous to redesign the tooth shape in trying to reduce tooth loss. Semicircular shapes seem to offer greater resistance to losses caused by sudden starting and stopping. 9.4.2

Conveyor belts

Conveyor belts are used to move solid (powdery or chunky) materials over distances which can range from a few metres to tens of kilometres, and are not limited to strictly horizontal transport but can accommodate slight elevation differences - the permissible conveyance angle depending on the type of material and the physical characteristics of the belt’s surface. In nearly all instances, the belts are prevented from sagging under the payload by rollers, placed at regular intervals under the working (loaded) portion and even under the idle (returning) part of the looped belt. The force needed to move the load is transmitted at one end of the loop by a driven roller while the idler roller’s main role on the other end of the loop is to provide the tension necessary for friction-generation at the driver roller and to minimize sagging in between the supporting (and idling) rollers. Continuous reinforcement, in the form of a parallel array of cables or cords made of steel or high modulus polymeric fibres (see Fig. 9.18 for some typical belt constructions), restrains the flexible matrix from elongating in the loop direction. Additional short fibre reinforcement in that direction (i.e. parallel to the movement of transport) would have only a very limited, if any, benefit.

236

Short fibre-polymer composites

conveyor belt reinforced with cable

additional fabric reinforcement

conveyor belt reinforced with c o r d (or fabric); fill not shown

9.18 Typical conveyor belt constructions.

Conveyor belts have to be reinforced perpendicular to the travel direction for two important reasons: (1)without such reinforcement, prevention of separation between the continuous reinforcing cables (known as ‘ripping’ or ‘slitting’)would be left to the relatively weak matrix; (2) under payload, the cross-section of the

Design and end-use applications of short fibre filled rubbers

237

belt would undergo side-to-side sagging between the supporting rollers, causing the cargo to shift constantly. To provide lateral stiffness, some belts are built with plies of woven fabric in which the warp provides the continuous, longitudinal reinforcement and the fill the side-to-side reinforcement. The use of fabrics alone is more common in belts of low physical demand. When the loads are appreciable and, therefore, the tensions high, the previously mentioned parallel cable/cord array is used longitudinally, overlaid (sometimes on both sides) with woven fabric reinforcement. The woven fabric, if placed close enough to the surface of the conveyor belt, also offers protection against ripping and gouging when jagged bulky loads are dumped onto the take-upend of the belt. Cutting can also occur when such a load shifts position on the belt, a situation quite common during elevated transport. These ‘lateral functions’ of the fabric can be improved by the use of short fibre usually in addition to the fabric - dispersed in the matrix in such a manner that the fibres are oriented side-to-side and that the fibres are concentrated, if possible, as near to the two surfaces of the belt as feasible. These criteria are obvious from the property improvement discussion (section 9.3). The orientation is necessary to reduce extension of the matrix across the width of the belt. The closeness to the top surface (next to load) is needed to improve cut resistance and to the bottom surface to increase bending stiffness for reduced side-to-side bending (known as ‘troughing’).The actual payload will determine what kind of reinforcement and in which part of the belt it is to be used. For instance, if the belt is to carry sharp rocks or ores, cut resistance will be the primary consideration. The amount of fibre in the matrix does not need to be excessive: with polyamides and polyesters, 5-lophrfibrecontent shows significant improvements. With high modulusfibres (e.g. aramids), 3-5 phr may result in much better performance. For carrying heavy payloads, conveyor belts with several layers of reinforcement (both lengthwise and crosswise) are used. Some of these constructions can be quite elaborate combinations of warp sheets and heavy fabrics separated by layers of rubber, the ‘interply’. The thickness of the interplies determines the amount of shear the reinforcing layers will generate while going around the end drums, etc. Harandi et a1 subjected a three-ply conveyor belt to rigorous mathematical structural analysis and found that to attain minimum force difference between adjacent plies the interply thickness should be equal to that of the fabric ply.30 If short fibre reinforced rubber is substituted for the conventional, non-reinforced interply, this conclusion should be scrutinized. There is a special consideration, unique to conveyor belts. While power transmission belts are manufactured as a fixed-length continuous loop, conveyor belts - which can sometimes extend several kilometres - are manufactured open-ended to be spliced together either into a longer piece or to form the loop. There are many means of splicing; metal clips, and hooks (staple-like fasteners) are used for light and medium duty conveyor belts. Heavy duty belts require extremely complicated layer-by-layer overlapping of the reinforcing cords, fabrics and cables along with additional thermoplastic matrix layers which are then vulcanized or thermoformed in situ using special heating equipment.

238

Short fibrepolymer composites

ner tube

spiral construction

braided construction

9.19 Some basic reinforced hose constructions.

Regardless of the type of splice used, the splice is always the most vulnerable area of the belt. It is always advisable to reinforce the matrix material around and within a splice with extra fabric but especially with fibre-reinforced material. The cost of the fibrc, even at relatively high loadings (e.g. 15 phr), is insignificant compared to the advantages gained. With the metal fasteners, the clips or hooks may tend to pull out of the matrix, even when a woven reinforcer is used. With heavy duty belts, the extra splicing and material build-up causes a thickening of the belt which, being stiffer than the rest of the belt, then undergoes more severe contortions at the end rollers. 9.4.3

Hoses

Hoses are designed to contain fluids, either to move them along the hose accompanied by pressure drop, or to transmit force with no or little pressure drop and with no or little movement of liquid. Except for very low pressure tubing applications, hoses have to be reinforced to avoid either bursting or radial and/or longitudinal expansion due to the internal pressure exerted by the liquid. When a hose is used to transmit liquid along its path, it is usually not vital to prevent a change in its shape, either lengthwise or radially. The reinforcement is needed mainly to increase the burst performance of the hase to meet the maximum anticipated pressure within. However, when the hose is used to transmit force either as a signal or as a power actuator (e.g. in brake lines), it is imperative that the volume of the hose changes as little as possible. Hoses in such applications are usually referred to as hydraulic hoses. In either case, the hose has to be reinforced with an array of continuous filamentary material: braid, woven fabric or spiralled cord/cable. Some basic hose constructions are shown in Fig. 9.19. A hose contains an inner tube, its composition chosen to withstand the chemical nature of the liquid (i.e. it has to be impervious to the liquid) for the lifetime of the hose. This tube is usually a rubbery elastomer or soft thermoplastic in low pressure applications, and a hard thermoplastic or thermoset for high pressure usage.

Design and end-use applications of short fibre filled rubbers

239

The inner tube is not designed to withstand pressure by itself. The pressure is counteracted by the reinforcing layer(s) tightly wound around the inner tube. The nature and number of layers of this reinforcement will determine the pressure capabilities and the dimensional stability of the hose. For instance, while braid reinforcement can be perfectly acceptable for hoses meant to convey liquids even at great pressures, it would be impossible to maintain volume consistency with such a construction, no matter how many layers are wound on the inner tube. Since the burst pressure is inversely proportional to the reinforcement’s diameter, it is preferable to use hoses with as small an overall diameter as possible; therefore, high modulus, high strength reinforcers are used in all but the inexpensive, bottom-of-the-line products. Nylon, polyester, glass, steel and aramid (in order of increasing moduli and specific strengths) have been used successfully as such continuous reinforcement. To maintain dimensional integrity,spirally and oppositely wound layers ofhigh modulus cords or cable are used in high pressure hydraulic hoses. The spiral angle is critical in obtaining constant dimensions both axially and radially. The reinforcing layer(s) are then covered with a layer of rubber or thermoplastic to protect the reinforcement from ambient damage, both physical and chemical, including ultraviolet rays. Short fibre reinforcement can be very successfully applied in two areas of the hose’s construction. The obvious one is the cover stock where improved cut and abrasion resistance are always needed. If the cover over the reinforced core is extruded, it must be remembered that the short fibres will align in the direction of the flow, i.e. they will be parallel to the length of the hose. This could add to the overall bending stiffness of the hose if too much fibre is used. A special technique (described next for the entire hose) could be used to obtain at least partial orientation along the hoop (circumferential) direction of the cover, thus avoiding contribution to bending stiffness. A special technique, developed by M ~ n s a n t o ,allows ~ the reinforcement of relatively low pressure hoses with short fibres alone, without continuous reinforcement. Since the hoses are extruded, a special mandrel is needed to allow choosing the fibre orientation from wholly axial to nearly completely circumferential. A very detailed description of this technique is given for an EPDM/ cellulose fibre combination by Goettler and SwiderskiZ8 Short fibre reinforcement has another, very viable, use in the reinforcing layers of a hose. Sometimes the pressure requirements arid the reinforcing fabric’s nature are such (e.g. with a braid or spirally wound cords not laid down tightly side-by-side) that there is an unreinforced gap through which the inner tube can be extruded under constant or overload pressures. Since all reinforcing layers are either encased in an elastomeric, very thin ‘sheet’ or are separated by such elastomeric sheets to prevent chafing on each other, fibre reinforced sheets may be used to bridge these ‘tensile gaps’. Sometimes, fibre reinforced sheet is wrapped on to the inner tube first to prevent a tightly wrapped first reinforcing layer from impinging on a relatively soft inner tube. Another application, often overlooked, is the reinforcement of hose ends with short fibre reinforced strips where the coupling is joined on to the ends of the hose.

240

Short fibre-polymer composites

The nature of these couplings is such that they sometimes bite into the hose, damaging its reinforcing layer(s). Application of additional fibre reinforced cover strips and use of slightly larger couplings may alleviate this problem. A rarely used application where short fibre reinforcement could be put to good use is in hoses operating under partial vacuum. Most of these contain an inner spiral; usually metallic, to prevent them from collapsing and the purpose of the wall material is basically to prevent air from leaking in along the hose’s length. If the sheeting (usually the wall is thin) were to be reinforced with short fibres, especially if oriented perpendicular to the spiral, the spacing of the spiral’s pitch could be made larger, saving on mass and cost. Whether large, kink-resistant vacuum hoses can be made without the spiral reinforcer has not been demonstrated yet. Considering that in order to kink the hose the wall has to bend, it is not impossible to image a vacuum hose reinforced solely by short fibres against collapsing. There is a gamut of rubbery, plastic, and thermoplastic materials used in To decide whether a hose failure could have been avoided by proper reinforcement, refer to SAE Technical Paper No. 91057.34 9.4.4

Friction products

The applications belonging to this category all depend on the ability of fibres added to a matrix either to increase its frictional properties or to provide mechanical integrity during its end use at elevated temperatures. Contrary to all previous applications where the success of fibre reinforcement depended on controlled fibre orientation, friction applications require randomly dispersed fibres in the matrix. Short fibres in the form ofasbestos have been used for a long time. Except for its pulmonary effect, asbestos is an excellent fibre for friction products: it can withstand extreme temperatures, has sufficient structural integrity and is inexpensive. There have been several attempts to replace asbestos in brakes. Steel fibrids have been used in the so-called semi-metallic brake shoes. The disadvantages of steel are rusting, low cold friction and high density. Glass fibres showed poor wear resistance in the pads, aggressive wear of the mating parts, and difficulty in maintaining the original fibre length during the preparative steps. The last problem (plus cost) also made the use of carbon fibres unattractive. Cellulosic fibres have too low strength and modulus, and also char at relatively low temperatures. Synthetic thermoplastic fibres have too low melting points to be considered for more than a single-occasion use. The only fibre so far which can match asbestos’s performance (except for price) has been p-aramid, either in the form of staple or pulp. Staple is obtained by cutting continuous length yarn bundles into - usually - equal-length pieces; the pulp is obtained by grinding the staple in a machine called the ‘refiner’. The pulp not only has random length and shorter fibres than a corresponding staple, but because the fibre is composed of axially highly splittable polymeric chains, it has numerous fibrillated ‘branches’ attached to the original trunk. These branches

Design and end-use applications of short fibre filled rubbers

24 I

are not only shorter but also much thinner than the trunk and have jagged surfaces, making them more surface active, both mechanically and physicochemically. The Du Pont Co with its Kevlar@aramid pulp pioneered the replacement of asbestos. Most of the early publications”-39 were published by its researchers. The same machinery and about the same preparative steps can be used in making the finished friction material as for asbestos. However, some modifications both in the machinery and in the technique are necessary when trying to disperse the fibres in the first steps.40 It is advantageous to use the so-called ‘masterbatch’ which is a preblend of p-aramid pulp, a particular elastomer and a particulate filler in about 40:40:20 weight ratio. More details on how to facilitate the dispersion of fibres can be found in Ref 4. Short aramid fibres have been used worldwide not only in automotive disc and drum brakes, in truck brake blocks, and railroad brake shoes, but also in industrial brake and clutch facings, and in manual (automotive) clutch facings.41 Since 1987, Volkswagen vehicles have been equipped with asbestos-free clutch facings containing aramid fibres.42 A very special application, ‘friction paper’, is described below. In clutch facings 50% asbestos can be replaced with 20% p-aramid; in disc brake pads, drum brake shoes, truck block brakes, and railroad brake shoes less than 5% p-aramid fibres are ~ u f f i c i e n t . ~ ~ 9.4.4. I

Friction paper Friction paper is a specialty application where a combination of fibres, fillers and elastomer is used as a precursor to the finished product which is basically a random fibre-filler mass saturated with phenolic resin. The friction paper, used in automatic transmissions, is the equivalent of the clutch plate in a standard (manual) transmission. While with a manual transmission the gradual engagement of the clutch depends on the skill of the driver, in an automatic transmission the otherwise quite instantaneous engagement is mitigated by the unique design of the clutch assembly. In very simple terms, this assembly consists of a series of thin friction paper discs separated by a set of thin metal discs, all rotatingfreely on a shaft until the two endplates are forced together rather abruptly. To reduce the jerking effect, the friction paper contains a large percentage (60% or more) of voids, all filled with transmission oil. At the moment ofengagement, first the thin layer of transmission oil, then the friction paper discs provide increasing torque as more and more oil is being pushed out of these voids with increasing compression. The purpose of the friction paper is to provide enough mechanical strength, good frictional properties, adequate but limited compressibility - all this at elevated temperatures, and sufficient void structure. The void structure of the fibres is achieved by using paper-making techniques (hence the name friction paper). To provide sufficient mechanical strength, the entangled fibre mass, which is randomly oriented in a plane, is temporarily reinforced with a small amount of elastomer, applied in a latex form. Various fillers are added, not only to reduce the cost but also to modify the frictional and wear characteristics. This

242

Short fibre-polymer composites

fibre-filler-elastomer composite is then transferred to the next processing stage where it is saturated with a liquid phenolic resin. Then the whole assembly is press-cured into thin sheets and onto metal backing discs. The fibres used include a combination ofcotton, asbestos and p-aramid pulp, the latter added to increase both mechanical strength and thermal resistance. The most commonly used reinforcing latex is NBR. 9.4.5 Rocket motor insulator

Rocket motor insulators, sometimes called liners, are that part of a solid fuel propelled rocket which protects the metal casing from the hot combustion gases generated as the propellant burns inside the casing. Originally, asbestos filled phenolic resins and elastomers were used for this Asbestos loaded nitrile butadiene rubber (NBR) was selected for the booster rockets of the space shuttle program; its success is attested by the fact that it is still being used in manned flights. The typical booster burns for about 2 min during the first-stage ascent and generates gases close to 3200 0C.4s The actual insulation is, however, provided not by the elastomer itself but by the char that forms on its surface as the temperature rises: during endothermic pyrolysis, volatile products evolve and a carbon residue is left behind. This char region advances throughout the insulator as the pyrolysis zone advances into the virgin material. However, only a limited thickness of char builds up because the crumbly outer layers rapidly erode due to mechanical shear forces (gas velocities reach more than 160m s-’), thermal stress and the internal pressures generated by the ~ o l a t i l e s To . ~ ~counteract the ablation of char, fibrous fillers are needed which can withstand both the mechanical forces and the high temperatures. Save for health and environmental concerns, asbestos is an ideal reinforcing fibre. However, the tightening ofexposure limits both in the State of Utah (where major rocket motor manufacturers are located) and by the US Federal Government prompted manufacturers to look for alternative fibres. Above Mach number 0.1 asbestos-phenolics are more resistant to erosion than all other candidates. Since the char layer is only weakly attached to the virgin insulation, it tends to be prone to mechanical shearing at high velocities. When the flow does not exceed Mach number 0.09, only thermochemical ablation takes place and p-aramid was found to be an excellent replacement for asbestos, especially in conjunction with EPDM rubber. This combination was recommended after asbestos-phenolic, Kevlar-EPDM, glass-phenolic and carbon fibre-EPDM items were compared for mechanical and bond strength, as well as ‘torch time’. The latter is the time required to reach a pre-selected temperature on the back, i.e. unheated, side of the sample when the flame of a ‘standard’ torch is applied to the front in a controlled manner. The p-aramid-EPDM combination was the only item which exceeded (by more than 30%) the torch time of the asbestos-phenolic item.46 A typical, simplified, EPDM composition is given below:

243

Design and end-use applications of short fibre filled rubbers

1.6

c

I v)

1.2

E E

CI

0.8

0.4

0 0

0.05

propellant gas velocity

0.1 0

0.1 5

[ave. Mach number)

9.20 Erosion of various composites versus gas velocity. A = 35 phr HiSil/ EPDM, B = lOOphr asbestos/polyisoprene, C = 65 phr HiSil, 20 phr Kevlar/polyisoprene, D = Kevlar/EPDM.

Nordel 1040 Neoprene FB Silica Antioxidants Kevlar, 1/4-in (0.64mm) staple Sulphur Accelerators

80 phr 20 36 3 16 I 2.25

When comparing the p-aramid-EPDM system with others, it was found that p-aramid-EPDM resulted in significant reduction in the erosion rate of the char (see Fig. 9.20). One significant advantage of EPDM-based insulators over the phenolics is that the EPDM compound is amenable to standard rubber manufacturing techniques. The primary prerequisites for a reinforcing fibre in ablative applications are good thermal resistance, substantial mechanical properties at very low and very high temperatures, low thermal conductivity, low specific gravity, high reinforcing efficiency, no physiological hazard and no smoke generation. Thermal resistance implies high or no melting point (p-aramid does not melt; it starts to

244

Short fibre-polymer composites

decompose at around 500 "C).Good residual mechanical properties are not only needed at high temperatures (during combustion) but also at very low (e.g. arctic or outer space) temperatures depending where the rockets are going to be stored prior to firing. Low thermal conductivity is needed so as not to counteract the thermal insulation provided by the char. Low specific gravity results in higher payloads. High reinforcing efficiency permits using less reinforcing fibre which results both in higher payloads and in reduced costs. The inhalational hazard, such as presented by asbestos during both manufacturing and firing, must be avoided. The smoke generation criterion is aimed specifically at military uses of rockets. In tactical military rockets not only the propellant but also the insulator (both the elastomer and its reinforcing fibre) have to be smokeless so as not to leave telltale signs of the launch location. An ancillary, ablation-related, application of fibre reinforced elastomers is the gasket material occasionally used with rocket motor casings. Because of its size, the rocket boosters of the space shuttle are assembled from four cylindrical segments. Each segment is manufactured by first coating the inside metal surface of the casing with an adhesive, then bonding the asbestos and silica-filled NBR to the casing and vulcanizing. Afterwards a liquid elastomer, also containing asbestos and silica, is applied - to provide the bond between the insulator and the propellant - and, finally, the propellant is cast in the required shape. These four segments are joined using tang-and-clevis field joints.4s Where these segments meet, gaps are left in the propellant and insulator which would allow access to the casing by the hot gases, causing eventual localized failure (such as happened with the Challenger). Huge O-rings are applied (currently three per joint) to prevent the combustion gases from reaching the casing. The general prerequisites mentioned above for the rocket insulator hold true for these gaskets too, with two additional requirements: high resilience and low compression set at low temperatures. 9.4.6 Automotive uses

These applications fall into two general categories: in tyre uses and body panel reinforcements. Seals used for automotive purposes are mentioned in the next section (9.4.7).The in-tyre usage can be separated into solid and pneumatic areas. Solid (industrial) tyres have been dealt with adequately by Goettler and Swiderski.28 They pointed out that highly oriented fibres in the circumferential direction in the base rubber reduce the dimensional instability (caused by sudden stop-and-go operations, such as in forklifts) of the tyre, hence its slippage on the rim. One disadvantage which has not been mentioned earlier is that, by adding short fibres to the base rubber, its heat build-up potential increases significantly. At elevated temperatures, the adhesive bond between the rim and the base rubber is severely compromised and leads to exactly the same failure, slippage, that the fibre reinforcement was supposed to remedy. Therefore, careful experimentation is needed to determine whether one's operation would actually benefit from the fibres.

Design and end-use applications of short fibre filled rubbers

245

Pneumatic tyres are built from many components (Fig. 9.21). The carcass ply (or plies), wrapped around the (usually steel) bead, provides the force and load transfer between the tread area and the tyre r i a . The carcass plies are protected from damage by the sidewall rubber and by belts and/or ‘undertread’ which separate it from the frictional surface, the tread. As the sidewall undergoes frequent and severe deflection, the area around the bead is subjected to shear, extension, and compression. To increase the stiffness (therefore moving the so-called hinge point from the rim flange area into the relatively thin sidewall where it generates less heat) and to protect the carcass from abrasion, the bead area is reinforced with an apex, wrapped in a chafer. The inner liner is the real barrier against air leakage. Nearly all these components lend themselves to short fibre reinforcement. The sidewall of a radial tyre, with its one or two radial cord plies oriented in the toroidal direction (i.e. cords running bead-to-bead in the plane, or nearly so, of the axis) and with the very thin rubber covering these plies, is the most vulnerable to physical damage. Cuts in the sidewall not only mean loss of air: the damage is usually irreparable. Cut resistance of the sidewall can be increased with little additional heat generation if fibres are incorporated near the outside of the sidewall rubber, circumferentially oriented, perpendicular to the toroidal plane. Chafers, as the name implies, were originally used to prevent chafing of the carcass cords by rubbing against the rim. These are basically thin, plain woven, rubber-impregnated fabric strips wrapped around the bead at a ‘bias’, i.e. that both the warp and fill are at 45 angle to the bead where they cross. Adding short fibres to the surface of the fabric could increase the chafing-resistance, but the thinness of the rubber matrix makes this approach difficult. However, replacing the woven chafer fabric with a fibre reinforced rubber strip is a viable alternative.28 The apex (a solid rubber piece with a roughly triangular shaped cross-section) has been successfully reinforced with short fibres to give it more stiffness, or rather greater bending resistance. As mentioned before, the apex (also called bead filler) is designed to provide a smooth spatial transition for the belt fabrics as they turn around the bead and to move the flex-zone farther away from the bead area. The width of the chafer fabric is chosen in such a way that its upturned edges do not coincide with the tip of the apex, nor with each other. This design spreads a potentially damaging major hinge point at the bead away to several mini-hinge points. In order for the apex to be properly reinforced, that is to increase its bending resistance, the fibres have to be oriented in the base-to-tip direction. Therefore, simple extrusion of a wedge-shaped strip with the reinforcing fibres in it is not going to be beneficial; the wedge either has to be built up in the mould from properly oriented strips and then cured, or one may attempt to use the expanding-die method demonstrated for hoses to obtain fibre orientation perpendicular to the extrusion direction. The inner liner material does not have to be strong, just impervious to air. However, during the curing process (especially with a tight ‘drumset’ and/or when tyre cords with high thermal shrinkage are used), the carcass cords tend to O

Short fibrepolymer composites

246

tread

-belts

- carcass plies

- bead

inner Ii

sidewall

er

filler

carcass turnup (b)

\bead filler

9.21 Typical components of a radial tyre.

Design and end-use applications of short fibre filled rubbers

247

pull into the inner liner. This usually manifests itself in its milder form as a visible indentation on the inside of the tyre; in its severe form, the carcass cords actually pull through the inner liner, thereby breaching its integrity. To avoid this, short fibre reinforcement, with the fibre direction perpendicular to the carcass cords, is w r y beneficial. Relatively small amounts, 5phr or less, of staple or floc (e.g. inexpensive nylon) can completely eliminate the pull-through. In some cases, even the thickness of the inner liner can be reduced without the threat of cord strike-through. Both abrasion and cut resistance could be improved by incorporating the fibres in the tread pattern but this is not practical with current manufacturing methods. The fibres would have to be perpendicular to the surface of the tread to be effective for improved tread wear and parallel to it for increased penetration resistance. No efficient method is available for producing the precursor of the tread (sometimes called the ‘camelback’) with such fibre orientations, nor would the orientation remain undisturbed during the moulding and curing process. ’ Belts d o not need fibre reinforcement nor is there any practical means of achieving it, especially since the rubber surrounding the belt cords is so thin. To increase the lateral integrity of the belt, it is essential to orient the fibres perpendicular to the belt cords. Otherwise, the fibres would cause additional heat generation without significant modulus improvement in the circumferential direction. Further information on tyre components is provided by Goettler and Swiderski.2 Prevorsek and co-workers4’ undertook a study to measure the heat generation rate of rubber reinforced with short nylon and polyester. They found, as expected, lower heat generation with the lower modulus fibre at fixed strain amplitude. What makes their paper interesting is that they developed a three dimensional finite element analysis technique. Very little has been reported on the use of short fibre reinforcement in the exterior panels of automobiles. General Motors did some theoretical investigations, reported by Kia.48 Glass was used as the reinforcing fibre. The main problem was that the fibres close to the surface left a visually perceptible mark, called the readout. The readout is basically deformation of the surface with a texture similar to that of the underlying fibres. The magnitude of the readout depends on the thickness of the fibre. With glass fibres of less than 10-20 pm in diameter, the readout did ‘not cause cosmetic problems’. 9.4.7 Miscellaneous end uses

Medium to light coated fabrics, including articles formed from them (e.g. diaphragms), can benefit from short fibre reinforcement. The objective is to reduce the end count of expensive twisted cords or strands in the fabric and substitute short fibre reinforcement in the sheeting material to close the gaps so generated. A combination of coated fabric and hose resulted in flexible couplings of low pressure, large diameter pipes, such as used in sewage lines. The flangeless cast

248

Short fibre-polymer composites

iron pipes are joined by rubber sleeves which are secured by stainless steel hose clamps on the respective pipe ends. These couplings - having to withstand some hydrostatic pressures and the resultant pull-out forces, as well as the lateral shear forces caused by pipe misalignment - are thick walled. Yet, during installation this thick sleeve has to be stretched over the pipe, an especially difficultjob in cold eath her.^' The sleeve could be made lighter, thinner and still easily manageable with short fibres incorporated circumferentially. The same kind of clamps could be used, but a size smaller. The use of fibre reinforcement in dock fenders and methods to fabricate them has been discussed by Goettler and Swiderski.28 This is an ideal application because the manufacturing method lends itself to easy implementation of the preferred fibre orientation, i.e. around the circumference of a tubular fender. Sheet roojing, which usually is an unvulcanized EPDM compound or similar, can benefit greatly from short fibre reinforcement. Such roofing material is generally applied to industrial buildings with large roof areas. Therefore, just the placement of large, long pieces of rubber sheeting can put considerable stress on the sheeting. In addition, any large unevenness in or on the roof will allow the unvulcanized material to creep, to the point that it will fail. Tensile reinforcement with short fibres is a practical solution to these problems. A similar but often overlooked solution is in green (1.e. uncured) rubber stock that is so weak that its work-up on the mill becomes impossible without reinforcement. Here the fibre reinforcement is just a temporary processing aid. Typical examples are: more than 10 times improvement in green yield strength of a Sol50 SBR-NR stock with 5 phr nylon or polyester staple of 6 mm l e r ~ g t h . ~ , ’ ~ Seals and gaskets are a potentially large market for short fibre reinforcement. The ‘border’ between seals and gaskets is somewhat blurred. Usually, seals are relatively soft, pliable, mainly elastomeric, have three dimensional, sometimes intricate shapes, and are used around moving parts. Gaskets are harder, contain less than 20% elastomeric ‘binder’ and are usually flat, being used between flat flanges, engine blocks, etc. Both are pre-manufactured off-site to the desired shape. (Sealants, discussed below, sometimes called formed-in-place gaskets, perform similar functions to gaskets but they are dispensed in a viscous liquid form at the site of use and harden later.) Seals and gaskets are not only important in fluid transfer and chemical/petroleum manufacturing, but also play a very important role in automotive engineering5’ and many other industries and uses too numerous to list. Their functions are also numerous: they have to seal, i.e. prevent fluids (liquids, gases and vapours) from leaking between two, not perfectly mating, surfaces; they have to allow motion between these surfaces whether caused by temperature variation or by vibration or by rotating shafts; they have to maintain resilience for long periods of time and under adverse ambient conditions; they have to distort under pressure to conform to the irregularities of the mating surfaces yet resist ‘extrusion’ under pressure; and they have to withstand chemical attacks and extreme temperatures. Fortunately, individual seals and gaskets do not have to meet all these requirements, just most of them.

-

Design and end-use applications of short fibre filled rubbers

249

What short fibre reinforcement offers to seals and gaskets is excellent creep resistance, especially under elevated temperatures (only with heat resistant fibres: p-aramid, glass, asbestos). Gaskets are made with two methods: compressed (C) gaskets which are calendered from a solvent-based mixture, and beater addition (BA) gaskets which are produced on machines similar to those of paper-making technology. Originally, asbestos was the main fibre choice for gasket reinforcement, but recently, health considerations have necessitated other choices. Cellulosic fibres provide sufficient strength and modulus for ambient temperature service but the fibres degrade at 200 "C. Glass has the thermal stability and the high strength/modulus but it tends to break up in high shear mixing, thus reducing its reinforcing efficiency. Beside the relatively high price, the disadvantage of p-aramid is its tendency for high static electricity build-up (a consideration to be taken very seriously when working with flammable solvents). Frances has presented a detailed paper on the replacement of asbestos in gasket application^.^' He found that among the three parameters used in judging gasket performance (tensile strength, measured according to ASTM F152; gas sealability, ASTM F37 Method B; creep relaxation or stress retention, modified ASTM F38), creep-relaxation was the most critical. He concluded that p-aramid pulp at greater than 5% levels far exceeded the sealability requirement ofboth BA and C asbestos gaskets; that about 8 % p-aramid pulp was needed to achieve the tensile strength (- 23 MPa) of the high performance C asbestos gaskets (BA asbestos gaskets are 13 MPa); that 8 % p-aramid was needed to equal the stress retention performance of C asbestos gaskets at 300 "C,and less than 3% to match the BA asbestos gasket performance. Based on these studies, he concluded that p-aramid could replace asbestos economically in the high performance, high price C gasket market, yet economically it could not compete with asbestos in BA gaskets. A manufacturingguide is available for using Kevlar" pulp in gasket ~heeting.'~ As mentioned above, seals are usually used on contact with moving, rotating parts. It is therefore important not only for the seal to exhibit high abrasion resistance, but also to impart little wear to its mate. Watson and ~ o - w o r k e r s ~ ~ * ~ ~ studied these effects in conjunction with short p-aramid fibres and found that 15 phr Kev1ar.Y.pulp in a neoprene compound reduced the coefficient of friction from 0.55 to 0.25 and increased the NBS abrasion resistance from 243% to 436% in the on-end orientation direction. In the ASTM D-3702 thrust washer test, 10 wt% lOmm long Kevlarg staple in a copolyester resin resulted in a weight loss of only 0.7 mg of the steel washer.53With thermoplastic nylon 6,6,20 wt% of 6mm Kevlar-49 staple produced a wear factor reduction from 917 to 239, a coefficient of friction decrease from 0.435 to 0.390, and a steel washer wear from 0.2 to <0.01 mg when run at 51 mm s--' and 1.72MPa.54 In bridge, motor and building mounts the purpose of rubber blocks is to allow limited movement or expansion of the item being supported owing to temperature changes or operational phases. The presence of short fibres in such blocks, oriented perpendicular to the force, is to decrease the tendency of the rubber -

-

-

-

-

-

250

Short fibre-polymer composites

under constant load even when ‘idling’ - to be extruded, squeezed out of its normal spatial confinement. This tendency to change shape with time even under non-extreme loads is typical of polymeric materials, and it is referred to as creep. Andrews, discussed the strain sensitivity effects in vibration isolators.55 There are two areas where short fibre reinforced elastomers offer performance benefits in athletic shoes:

1 On-end oriented fibres in the sole provide improved abrasion resistance and, thus, longer wear life or more uniform wear. In properly designed soles, elastomeric parts with high fibre concentration should be used in areas of more pronounced wear and according to the intended end use (e.g. walking versus basketball shoes). 2 According to Adidas, recent application of biomedical factors to athletic footwear resulted in the development of a torsion bar placed longitudinally across the instep area.56The torsion bar allows the forefoot and the heel to rotate independently while providing lateral stability by maintaining heel-totoe alignment. By altering the amount and placement of the fibres (Adidas uses Kevlarw) in an injection moulded thermoplastic material, the level and configuration of torsion can be modified. Solid rubber tank pads undergo very severe punishment: rough surface conditions and extremely high footprint pressures. The reinforcement of tank pads has been the goal of many studies. Limited advances have been made so far. The major obstacle is the extreme heat build-up, regardless of whether the pad is reinforced with short fibres, longer filamentary material or fabrics. Hot rubber is very susceptible to gouging and chunking. Another new, quite esoteric application involves sound deadening or sound insulation. A recent technical news item5’ describes the use of non-woven lead fibre sheets (e.g. curtains), sometimes coated with PVC, as an effective noise control material especially suitable to isolated noisy operations in factories and in the building trade. This principle could easily be adapted for elastomeric sheeting where, beside the sound insulation, other requirements such as flame retardancy (e.g. with neoprene) or aesthetics (e.g. drape or colour) are desirable. Toray Industries (Japan) manufactures the carbon fibre sheeting. As a potential outgrowth of this concept, one may visualize the use of lead fibre reinforced elastomeric engine blocks to dampen vibration and to reduce sound transmission from the engine to the body panels. A composite of nylon thermoplastic resin and short aramid fibres is now marketed5’ as a low-friction, low-wear material for applications where frequent rubbing is encountered, such as in gears, clutches, similar machine parts, and electrical switches. The composite contains 20% by weight Kevlar 5 aramid short fibres in Zytel nylon resin, and is said to be fivetimes more wear resistant than the nylon resin itself. Yet, this composite - which retains sufficient strength to be used in continuous operation at temperatures up to 130°C - poses little abrasive threat to its mating counterparts. Articles made with this kind of composite can be used to replace metal counterparts with considerable savings

Design and end-use applications of short fibre filled rubbers

25 I

800

-E CI

c

I

600

I

400

200

0 0

10

20

carbon fibre loading 9.22

30

40

(phr)

Effect of fibre loading on conductivity, measured at 27°C with thermally vulcanized, unfilled neoprene.

in weight. An added, very cost effective, advantage over metal counterparts is the fact that these resin-fibre articles can be produced by high volume injection moulding. Additional applications claimed59 are hard roll covers, oil well packings, bearings and bushings. De and co-workers60*6 investigated the potential of using carbon fibres in neoprene to shield against electromagnetic interference (EMI) and found that 3C40 phr carbon fibre loading was sufficient to make the composite a potential EM1 shielding material in the electronics industry. They also found that conductivity dramatically increased with added fibre content (Fig. 9.22). 9.4.7. I

Viscosity enhancement

What sealants, adhesives, coatings and paints have in common is the need for high viscosity during the solidification process to prevent sag and drip, yet low viscosity during the actual deployment for ease of application. There are two similar, yet different, physical phenomena which many times are referred to interchangeably:62(i) pseudoplasticity implies that the viscosity of the material is dependent on shear alone (the thick material thins with increasing shear and re-thickens as shear is reduced) but it is independent of time; and (ii) the term thixotropy (the name commonly used to describe either phenomenon) implies

252

Short fibre-polymer composites Table 9.1. Comparison of Kevlar” and PulplusTMfibrous thixotropes

Chemical form Maximum usable temperature (“C) Minimum usable temperature (”C) Average fibre length (mm) Fibre diameter (pn) Aspect ratio Surface area (m2kg-’) Specific gravity Tensile strength (MPa) Colour Susceptible to Moisture content (YO) at room temperature and 25% relative humidity Price

Kevlar”

PulplusTM

p-aramid 400 Cryogenic 0.8, 2 or 4 12 67, 167 or 333 800@ 10,000 1.44 3620 Yellow Hot, concentrated acids and bases

Polyethylene 120 - 70 1 1-20 1000-50 3000-4000 402 White Strong oxidizing acids and some chloro-hydrocarbons

5-6 Substantial

0.1 Inexpensive

0.91-0.97

that the material thins with increasing time of shearing; there is a time-dependent delay in the thickening process when shear is removed if the material is thixotropic. In the sealant and adhesive industry, the primary prerequisite is a pseudoviscous material which is thin under high shear (i.e. it is easy to extrude, pump or spray) and viscous at low shear for sag or slump resistance. True thixotropy is needed with paints where high shear thinning allows brushing, roller or spray application, while the time delay allows the brush and roller marks or spray droplets to coalesce into a smooth film; the subsequently developing high viscosity then prevents gravity from manifesting itself in drips and sags. In the past, particulate, powdery fillers (e.g. chalk or fumed silica) or asbestos fibres were added to the low viscosity material. During overprocessing, asbestos fibres break down; fumed silica, which enhances viscosity because of its hydrogen bonded network, may not recover this network, which is temporarily broken up during high shear processing. In addition, the health hazards associated with asbestos necessitated the search for replacement fibres. These new fibres, in addition to providing thixotropy, also improve the tensile characteristics of the matrices. The effects obtainable were demonstrated with two fibrous thixotropes,63 p-aramid pulp (e.g. Kevlaro) and a special polyethylene pulp (PulplusTM).They both thicken and build up viscosity through the random orientation and physical entanglement of their fibres and fibrils. Under shear, these fibres/fibrils align and the material thins. Once the shear is removed, the fibres disorient and the material returns to its original viscosity. By varying the pulp concentration, a broad range of viscosity control can be achieved which is unaffected by processing or ageing, resulting in consistent performance and increased shelf life. The basic differences of these two pulps are summarized in Table 9.1.

Design and end-use applications of short fibre filled rubbers

253

Tabk 9.2. Comparison of fibre reinforcers through slump test data Material

Time (sf’

Material

Time (sf’

Pure epoxy with 5% fumed silica

8 49

Neopreneb with 30 phr asbestos

82

with 0.4% Kevlara pulp

47

with 0.75% PulplusTM Pulp

48

with 7.5 phr Kevlar@ Pulp with 11 phr PulplusTM Pulp

102

84.

Material

Time (sf’

Vinyl ester 3.3 with 2% 10.3 hydrophobic fumed silica with 2% hydrophilic 5.8 fumed silica with 0.2% Kevlar@ 11.0 Pulp with 0.4% PulplusTM 11.7 Pulp

Needed to obtain a given length of slump. As solvent base adhesive.

The thixotropic effects in an epoxy matrix of these two pulps versus fumed silica indicate (Fig. 9.23) that 0.4 wt% p-aramid or 0.75 wt% polyethylene pulp can replace 5 wt% fumed silica without any adverse effects; the use of 1.5 wt% p-aramid pulp increased the viscosity at gravitational shear rate about 10-fold.In a polyester resin matrix, 0.2 wt% p-aramid pulp, 9.4 wt% polyethylene pulp and 1.5 wt% fumed silica resulted in practically identical viscosity versus shear rate curves. These three items were also compared in a variety of other matrices,63and results of various ‘slump’ tests (where gravitational sag was used as the source of low shear) are compared in Table 9.2. Because of slight differences in test methods, only data within a column are comparable. Because of both thixotropic enhancement and tensile improvement, fibre reinforcement can be utilized in a multitude of applications, e.g. in situ formed gaskets, sealants and adhesives (such as PVC plastisol adhesive and in silicone ~ e a l a n t ~highway ~), line paints, floor coatings, marine coatings, textured decorative coatings, roof mastics and other asphalt coatings; for roof coatings and other roofing materials see Ref 65.

9.5

Summary

Incorporation of short fibre reinforcement into elastomeric articles can be a powerful tool to impart all kinds of property enhancement, sometimes in a manner which would be impossible with conventional, particulate reinforcers. Indeed, the ability to provide directional reinforcement (anisotropy) is one forte of short fibre reinforcement. In most practical applications, the ability of the oriented fibres to counteract extension along their axes is put to use. Notable exceptions are friction products (where random fibre distribution is preferred) and ablative applications (where the thermal resistance of the fibres is of primary importance). When designing an article, fibre reinforcement throughout the entire body is usually neither needed nor is it always beneficial. Careful engineering and full understanding of the fibre’s properties and of the end use

254

Short fibrP-polymer composites

1o4

-C L

I

E to

103

z

Y

) .

.-cto. 0 0

.-to >

1o2

10 1o-2

lo-'

1

shear rate

10

1o2

103

(s")

9.23 Viscosity as a function ofshear rate with various fillers at 25 "C. Epoxy resin matrix (E) with A = 1.5wt% KevlarQ pulp, B = 4.9 wt% fumed silica, C = 0.4 wt% Kevlarm pulp or 0.75 wt% PulplusTMpolyethylene pulp; gs = gravitational sag, a-t = agitated tank, c-p = centrifugal pump, b = brushing. The shear rate for spraying is out of range and may be as high as lo5 s - ' .

functions of the article are needed for the cost- and performance-efficient utilization of short fibres.

References Abrate S, Rubber Chem Technol, 59 (1986) 384. Goettler LA and Shen KS, Rubber Chem Technol, 56 (1983) 619. Campbell J M, Prog Rubber Technol, 41 (1978) 43. Foldi A P, in Composite Applications, The Role of Matrix, Fiber, and Interface, Eds T L Vigo and B J Kinzig, VCH Publishers, New York (1992). Cheremisoff N P Ed, Elastomer Technology Handbook, CRC Press, London (1993). Goettler LA, Leib R 1 and Lambright A J, Rubber Chem Technof,52 (1979) 838. Coran A Y, Hamed P and Goettler LA, Rubber Chem Technol, 49 (1976) 1167. Uchiyama Y, Hosokawa M, Wada N and Ohino M, J Appl Polym Sci: Appl Polym Symp, 50 (1992) 283. Derringer GC, Rubber World, 165 (1971) 45. 10 Marker L F, US Patent 3,570,575 (1971).

Design and end-use applications of short fibre filled rubbers

255

11 Coran AY, Boustany K and Hamed P, Rubber Chem Technol, 47 (1974) 396. 12 Tanner D, Fitzgerald J A, Riewald P G and Knoff W F, ‘Aramid structure/property relationships and their role in applications developmznt’ in Handbook of Fiber Science and Technofogy,Vol 111, High Technology Fibers, Part B, Eds M Lewin and J Preston, Marcel Dekker, New York (1989). 13 Kevlar” UltrathixTMa New Thixotrope Specially Designed for Easy Dispersion, company publication H-29607, Du Pont Co, Wilrnington, DE 19898. 14 Morgan R A, ‘Reinforcement of Elastomers with Fluoroplastic Additives’, ACS Rubber Division meeting in Washington, DC, USA October 9-12 (1990). 15 Arnold C A, Hergenrother P M and McGrath J E, Composite Applications, The Role of Matrix, Fiber, and Interface, Eds T L Vigo and B J Kinzig, VCH Publishers, New York (1992) 3. 16 Pratte J F, Krueger W H and Chang I Y, High Performance Thermoplastic Composites with Poly(Ether Ketone) Matrix, Bulletin H-14471, Du Pont Co, Composites Div, Wilrnington, DE 19898. 17 Akhtar S, De P P and De S K, J Appl Polym Sci, 32 (1986) 5123. 18 Ramesh P and De S K, J Elastom Plast, 25 (1993) 106. 19 Foldi A P, Rubber Chem Technol, 49 (1976) 379. 20 Foldi A P, papers presented at ACS Rubber Division meeting in New York (1986)and at the International Rubber Conferences in Sydney (1988) and in Paris (1990). 21 Rubber World, 196(4) (1987) 45. 22 Outzs L L, ‘The use of aramid fiber pulp,/neoprene masterbatch to develop unique engineering and dynamic properties in rubber’, ACS Rubber Division meeting in Washington, DC, October 9-12 (1990). 23 Thomas A G and Southern E, Plast Rubb: Mater Applics, 3 (1978) 133. 24 Schallamach A, Trans lnst Rubber Ind, 28 (1952) 256. 25 Gayer M D, Waugh D L and Fisher D G Eds, Dayco Automotive Belt Drive Design Handbook, Dayco Corp, Dayton, O H (1980). 26 Kingley P Ed, 1987 Power Transmission Handbook, Penton Publishing, Cleveland, O H (1987). 27 Donaldson W K, Chapter 9 in Textile Reinforcement of Elastomers, Eds W C Wake and D V Wooton, Applied Science Publishers, London (1982) 225. 28 Goettler L A and Swiderski Z , in Composite Applications, The Role of Matrix, Fiber and Interface, Eds T L Vigo and B J Kinzig, VCH Publishers, New York (1992) 333. 29 Roger5 J, Rubber World, 143(6) (1961) 97. 30 Harandi J N, Kavousian A, Parvini M and Soufian M, Rubber Chem Technol, 65 (1992) 697. 31 Goettler LA and Lambright A J, US Patents 4,056,591 and 4,057,610 (to Monsanto) (1977). 32 R M A Hose Handbook, published periodically by the Rubber Manufacturers Association, Washington, DC. 33 Evans C W, Hose Technofogy, Applied Science Publishers, London (1979). 34 Miller S N, Systematic Techniques of Hose Failure Mode Analysis, SAE Technical Paper Series 9 1057, The Engineering Society For Advancing Mobility, Warrendale, PA 15096 (1991). 35 Loken HY, SAE Technical Paper Series 800667, The Engineering Society For Advancing Mobility, Warrendale, P A 15096 (1980). 36 Hoiness D E and Frances A, International SAMPE Symposium, Vol 32, SAMPE, Covina, CA 91722 (1987) 1173.

256

Short fibre-polymer composites

37 Space Age Fiber f o r Long-lasting Truck Brake Blocks, company publication E-89558, D u Pont Co, Wilmington, D E 19898. 38 When Stopping Really Counts . . ., company publication E-69531, Du Pont Co, Wilmington, D E 19898. 39 'Breakthrough in brake performance', DuPont Magazine, JanlFeb (1985). 40 Guide to Mixing and Processing KevlarB Aramid Fiber and Pulp, company publications E-53161 and E-65333, D u Pont Co, Wilmington, D E 19898. 41 Kevlar - A Reinforcing Fiber Substitutefor Asbestos, company publication E-38531, D u Pont Co, Wilmington, D E 19898. 42 Case History H-13657, D u Pont Co, Wilmington, D E 19898. 43 Galli E, Plastics Compounding, Nou/Dec (1981) 21. 44 Yezzi C A and Moore B B, 'Characterization of Kevlar/EPDM rubbers for use as rocket motor case insulators', AIAA/ASME/SAE/ASEE 22nd Joint Propulsion Conference, June 16-18, Huntsville, AL (1986). (Reprints available from American Institute of Aeronautics and Astronautics, 1633 Broadway, New York, NY 10019, USA.) 45 Seltzer R J, Chem Eng News, I9 Sept (1988) 7. 46 Woods C D, Myers F F Jr, Davidson T F and Ludlow T L, Replacement ofAsbestos in Smokeless Insulatorsfor Rocket Motors, Report No. M M T 38 11051 FR2, April (1984). (Available from Atlantic Research Corporation.) 47 Kwon Y D, Beringer C W and Prevorsek D C , Amer. Chem. SOC.Rubb. Div. Akron Group, Fall Meeting, Technical Papers (1988) 31. 48 Kia H G, Composite Applications, The Role of Matrix, Fiber and Interface, Eds T L Vigo and B J Kinzig, VCH Publishers, New York (1992) 295. 49 Cankovic M, Rubber Plast News, 26 October (1992) 57. 50 Mahishi J M, Rubber World, 207(2) (1992) 20. 51 Frances A, 1984 Nonwovens, TAPPI Symposium Notes, TAPPI, Atlanta, G A (1984) 125. 52 Manufacturing Guide for Gasket Sheeting Reinforced with K E VLAR'Y' Aramid Pulp, company publication E-69541 (2/91), Du Pont Co, Wilmington, D E 19898. 53 Watson K R and Frances A, Compos Polym, RAPRA reprint ISSN 0952 6919 (1988) 187. 54 Watson K R, Wu Y T and Riewald P G, Plast Eng, XLV (1989) 83. 55 Andrews F J, Rubber World, 207(2) (1992) 24. 56 Du Pont Magazine, N o 4 (1990) 1. 57 Techtextil-Telegrnmm, 22 March, No 27, Issue E (1993) 4. 58 Techte.util-Telegramm, 22 March, No 27, Issue E (1993) 3. 59 When Your Elastomeric Application Calls For Reinforcement. . .,company publication H-20485 (10/89) Du Pont Co, Wilmington, DE 19898. 60 Jana P B, Mallick A K and De S K, Composites, 22 (1991) 451. 61 Jana P B, Chaudhuri S, Pal A K and De S K, 'Electrical conductivity of short carbon fibre reinforced polychloroprene rubber and mechanism of conduction', Polym Eng Sci, 32 (1992) 448. 62 Frances A and Dottore J, Adhesives Age, 3 (1988) 27. 63 Through Thick and Thin, company publication H-24225, D u Pont Co, Wilmington, D E 19898. 64 Thixotropy with Kevlar," Arnmid Pulp, Bulletin E-91728, Du Pont Co, Wilmington, D E 19898. 6 5 Bulletin E-98171, Du Pont Co, Wilmington, DE 19898.

Index

ablation, 242, 243 abrasion resistance, 213, 226, 239, 247 acrylonitrile butadiene styrene (ABS), 3, 149, 166 ageing, 35, 39, 44, 45 aluminium flakes, 146, 147, 154-7, 163-6 anisotropy, 2, 5, 11-17, 23, 39, 40, 42, 60, 63, 87, 93, 94, 97, 105, 106, 112, 151, 169, 170, 193, 195-8, 202-8 applications, 1, 3, 18, 21, 22, 54-61, 71 aspect ratio, 11, 15, 45, 87, 90, 150-3, 159, 169, 170, 176-9, 185-7, 199, 200, 212, 214, 215, 217, 222 athletic shoes, 250 automotive body panel, 247 Banbury internal mixer, 91, 96, 193, 212 belt, 227-38 birefringence, 27 blends, 86, 88, 89, 97-9, 105, 106, 175 block copolymers, 85, 100 bonding agent, 117 brake materials, 240, 241 bridge mounts, 249, 250 building mounts, 249, 250 carbon black, 168-178 carboxylated nitrile rubber (XNBR), 91, 217 carrier mobility, 172, 181, 182 chemical resistance, 60, 166, 216 chlorobutyl NR, 222 clutches, 227, 241 compression blow moulding, 164, 165 compression moulding, 6, 61, 67, 88, 163, 164, 194, 199 compressive fatigue, 141, 142 compressive properties, 221, 224 contact microradiography, 17 continuous fibre reinforcement, I , 200, 201, 239 conveyor belts, 105,235-238

coupling agents, 24, 65, 72, 90, 158 creep, creep resistance, 21, 40, 41, 54, 64, 87, 248, 249 critical concentration, 149, 168 critical fibre length, 90 critical stretch ratio, 198, 202, 203 crystal nucleation, 31, 32, 39 crystal orientation distribution, 27 curing, 54, 61, 64-8, 73-7, 117, 194, 216, 218, 219; see also vulcanization clamping coefficient, 192 deformation mechanisms, 196, 197, 200, 207 density, 17, 36, 37, 44, 45, 60, 80 die swell, 100 clock fenders, 248 dough moulding compound (DMC), 2, 5, 17, 57, 69, 70, 75-81 dynamic fatigue, 135-139 dynamic mechanical properties, 90, 105-7, 122, 139. 192 electrical properties conductivity, 60, 144, 168-73, 175 effect of pressure, 176, 177 effect of strain, 177 dielectric properties, 60 resistivity, 144, 145, 148-50, 153-5, 170, 172-5 effect of temperature, 177-9 electromagnetic interference (EMI) shielding, 1447, 154, 165, 168, 182-8, 251 electron hopping, electron tunnelling, 150, 155, 170, 171, 172, 175, 178, 182 environmental attack, 42, 43, 166 ECPDM (ethylene-propylene diene modified), 86, 90, 91, 95, 176, 216, 239, 242, 243, 248 epoxies, 3, 54, 57, 58, 63, 70, 169, 174, 253 fatigue, 42, 90

258

index

fibre breakage, 4, 5, 23, 25, 34, 65, 66, 70, 71, 91.96.97, 110, 117, 118, 199, 214 fibre clumping, 67, 75 fibre length effects, 119-27, 130, 131 fibre-matrix adhesion, 2, 6, 7, 24, 41, 42, 57, 87, 89, 90, 126, 129, 156, 157, 196, 201, 220 fibrematrix separation, 220 fibre-matrix stress transfer, 8-1 1, 87, 90, 91, 101, 217 fibre orientation effect on properties, 87, 101, 102, 105, 195-208, 221-7 fibre orientation distribution, 2, 5, 26-33, 45-7, 66-70. 93, 94, 96, 105, 150, 176, 196, 213 measurement, 14-17, 29-31 modelling, 5, 6, 67 fibre properties, 17 fibre segregation (during moulding), 29-31, 8 1 fibre volume fraction. 5, 122, 126-33. 195-8, 205-8 fibres, 2, 148, 214-16 aluminium, 146, 158, 161, 165 aramid, 88, 214, 215, 220, 223, 224, 227, 229, 237, 24CL-243, 249, 252, 253; srr also fibres: KevlarTM asbestos, 214, 215. 240, 242, 244, 249 boron, 2, 96 brass. 165 carbon. 2. 23, 32. 88, 89, 91. 96, 98-100, 105, 106. 117, 158, 168-70, 172, 176, 177, 181, 182, 185-8, 215, 240, 242 cellulose, 3, 89-91, 95, 96, 106, 108, 109, 215, 232. 239, 240 copper, 165 cotton, 215, 221. 232, 242 glass, 2, 23, 64-66, 87, 91, 96, 117, 158, 161, 184, 214, 215, 221, 229, 240, 249 iron, 165 jute, 92, 96, 106, 215 KevlarTM,2, 23, 91-113, 212, 215, 249, 150 nickel, 174 nylon, 87, 89, 91, 94-6. 103, 11643, 215, 221, 222, 247, 248 polyester. 89, 91. 96, 215. 221, 232, 247, 248; see also fibres: poly(ethy1ene terephthalate) polyethylene, 215, 252, 253 poly(ethy1ene terephthalate) (PET), 87-91, 94, 95, 103, 105, 116-43 silk, 3, 87, 91, 96-8, 101, 215 stainless steel, 3, 146, 147, 158, 163-6, 229 fibrillation, 214, 240, 241 fibrils, 215 flame retardancy, I 1 I, I12 flat belts, 227-9, 234 flow of thermosets, 67-70 fracture, 41, 42, 103, 105, 1 I , 140, 141, 196-201 friction, 213, 225, 231, 241, 249

friction paper, 241, 242 friction products, 24&2 galvanomagnetic properties, 179, 181, 182 gaskets, 248 gears, 227 graft copolymers, 85, 86 Hall effect, 179, 181, 182 hand laminating, 63, 64 heat exchangers, 145, 146 high density polyethylene, 149 hoses, 95, 106, 23840 hydraulic hoses, 238 hysteresial heating, 136-9 image analysis, 14-17 impact properties, 41, 42, 59, 60, 71, 74, 87, 105, 156, 160-2 injection/compression moulding, 62, 63, 69 injection moulding, injection mouldings, 6, 11, 12-14, 21-41, 44-50, 57-63, 68, 69, 75-7.87, 15&2, 154, 155, 163, 164, 250, 251 intensive mixers, 96 interphase, 108, 129 isoprene rubber, 90, 216 limiting oxygen index (LOI), 11 I, I12 liquid crystal polymers composites, 3, 22, 39 long fibre composites, 23, 34, 41, 42 loss modulus, 122, 123, 132, 135 loss tangent, 106, 107, 123-5, 129, 1 3 2 4 temperature dependence, 124, 125, 129 low profile additive (LPA), 55, 56, 71 masterbatch, 212, 241 maximum packing fraction, 154, 156 mechanics of composites, 7-14, 90, 126-8 mica, 23, 24, 32 microwave anechoic chambers, 182 milling, 6, 88, 91, 94, 96, 97, 117, 121, 172, 173, 185-7, 193, 194, 199, 212, 213, 218 mixing, 4, 5, 25, 63, 88, 89, 91, 93, 96-8, 110, 117, 121, 162, 163, 172, 173, 179, 180, 185-7, 193, 199, 205-7, 212, 216 mixing index, 205-7 moisture absorption, 60, 74 molecular orientation, 27, 44, 45, 59 Mooney-Rivlin equation (modified), 195, 197, 198 Mooney viscosity, 218, 219 morphology, 5, 25-35 motor mounts, 249, 250 multiple V-belt, 230, 234 natural rubber (NR), 86-91, 94,95, 97-101, 105, 106, 109, 216, 221, 224, 227, 248 neoprene, 232, 253; see also polychloroprene

Index network, 144, 145, 149-51, 168, 169, 173-6, 178, 195 nickel flakes, 165 nitrile-butadiene rubber (NBR), 86, 91, 100, 106, 176, 242 noise: damping, 192 nylon composites, 3, 17, 22, 28, 29-31, 41, 42, 45-8,88, 149, 158, 166, 237, 250 nylon fibre reinforcement, 1 1 6 4 3 optimum fibre concentration, 95 paracrystallinity, 85 particulate reinforcement, 23 penetration resistance, 224-8, 247 percolation theory, percolation threshold, 149-55, 164, 168-70, 176 phenolic composites, 3, 54, 59, 64, 71, 72, 75-80, 149, 150, 242 pipe couplings, 248 polyamide composites see nylon composites polyamide thermoplastic elastomers, 85 polyaniline, 144 polycarbonate (PC) composites, 3, 22, 149, 158 polychloroprene (and composites), 89, 100, 116-43, 193-208, 216 poly(dimethy1 siloxane), 86 polyester composites, 3, 54-7, 63, 64, 70, 71, 75-81, 184, 217, 237, 253 polyester thermoplastic elastomers, 85 polyether thermoplastic elastomers, 85 poly(ether ether ketone) (PEEK) composites, 3, 22 poly(ether ketone) (PEK) composites, 22 poly(ether imide), 149 poly(ether sulphone) composites, 3, 22 polyethylene, 86, 98-101, 105, 106, 149, 165, 173, 175, 217 poly(ethy1ene terephthalate) composites, 22, 43, 166, 217 poly(ethy1ene terephthalate) fibre reinforcement, 11643 polymide composites, 3 polyisoprene, 90, 243 poly(methy1 methacrylate) (PMMA), 71, 86 poly(pheny1ene oxide)/polystyrene alloy (PPO/PS), 149, 166 poly(pheny1ene sulphide), 149, 166 polypropylene composites, 3, 17, 22-4, 29, 324,4G3,49, 149, 158-62, 184, 185 polypyrrole, 144, 185 polystyrene composites, 32, 185 polystyrene thermoplastic elastomers, 85 polysulphone, 149 polyurethane, 63, 217; see also thermoplastic polyurethane elastomers poly(viny1 acetate), 56, 65, 71 poly(viny1alcohol), 158

259

poly(viny1chloride) (PVC) composites, 42, 166, 170, 217 positive temperature coefficient composites, 153 power transmission belts, 135, 227-35 precision drive belts, 234 pressure-sensitive conductivity see electrical properties: conductivity properties of short fibre composites, 17, 18 pseudoplastic, 100, 251, 252 pulp, 215, 219,220- 227,240,249,252, 253 pultrusion, 163 radar absorbing materials, 182 reaction injection moulding (RIM), 63 reinforcement effectiveness, 199, 205-7 residual stress, 27, 42, 48-50, 55, 65, 71, 73-7 RFL (bonding agent), 117, 128, 140 rocket motor insulation, 242-4 roofing sheet, 248 rubber, 3, 21G54; see also natural rubber

SCORIMmprocess, 6, 34, 47 seals, 93, 248 self-adhesion behaviour, 109, 110 sepiolite, 3, 23 shear modulus, 195, 202-5, 214 sheet moulding compound (SMC), 2, 5, 6, 17, 67, 166 shielding effectiveness, 183-8 shrinkage, 35-40,45, 55, 56, 71, 77 silver particle filler, 149, 150 size, 2, 65, 72 solvent resistance, 216 sound insulation, 250 spherulites, 26, 27 splicing, 237, 238 staple, 215, 219, 220, 222, 227, 240, 247, 248 storage modulus, 90, 105-7, 122-8 strain-induced crystallization, 195 stress relaxation, 106, 108, 140-2 stress-strain behaviour, 1014, 195-203, 220-3 styrene-butadiene rubber (SBR), 89,91, 100, 174, 216, 221, 248 styre1.e-butadiene-styrene (SBS), 85 styrene-isoprene-styrene (SIS), 85, 89, 100, 101, 105 superconductors, 144 swelling behaviour, 118-20. 140 synchronous belts, 234 tack, 218, 219 talc, 23, 24, 32, 184 tank pad. 250 tear resistance, tear strength, 103, 105, 225, 226 temperature resistance, 24, 54, 60, 165, 216, 217, 249

260

Index

tensile strength, 17,90,95-98,1OC-3, 146,147,

149,15660,249 thermal properties, 42,145,156,157,166, 243 thermal conductivity, 42,48,59,60,145, 146,156,157,178,243,244 thermal expansion, 42,50,178 thermoplastic elastomers (TPEs), 84-1 15 thermoplastic polyurethane (TPU) elastomers,

85,88,91-113 thixotropic, 251,252 timing belts, 227,234,235 transcrystallization, 31 tyres, 95,105,106,216,2447

vacuum hose, 240 vibration: damping, 192,250 vulcanization. 173-5 warping, 21,24,35,39,71,77 water absorption, 131-5 wear, 213,232,234,241,247,249,250 weathering, 42,43,49 weld lines, 33-5,63,165 whiskers, 215 X-ray microradiography, 29,32;see also contact microradiography Young’s modulus, 11-14,17,448,64,67,

V-belts, 93,105,106,216,228-35

77-9,87,102,103