Polymer Composites: Proceedings, 28th Microsymposium on Macromolecules, Prague, Czechoslovakia, July 8–11, 1985 9783110856934, 9783110109948

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Polymer Composites

Polymer Composites Proceedings 28th Microsymposium on Macromolecules Prague, Czechoslovakia, July 8 -11,1985 Editor Blahoslav Sedlâcek

W G DE

Walter de Gruyter • Berlin • New York 1986

Editor Blahoslav Sedldcek, PhD., DSc. Institute of Macromolecular Chemistry Czechoslovak Academy of Sciences Heyrovsky sq. 2 CS-162 06 Prague 616 Czechoslovakia

Library of Congress Cataloging in Publication Data Prague IUPAC Microsymposium on Macromolecules (28th : 1985) Polymer composites : proceedings / 28th Microsymposium on Macromolecules, Prague, Czechoslovakia, July 8-11,1985. Includes bibliographies and indexes. 1. Polymeric composites-Congresses. 2. Polymers and polymerization-Congresses. I. Sedlacek, B. (Blahoslav) II. Title. TA418.9.C6P726 1985 620.1'92 86-6358 ISBN 0-89925-203-6 (U.S.)

CIP-Kurztitelaufnahme der Deutschen

Bibliothek

Polymer composites : proceedings / 28th Microsymposium on Macromolecules, Prague, Czechoslovakia, July 8-11,1985. Ed. Blahoslav Sedlacek. - Berlin ; New York : de Gruyter, 1986. ISBN 3-11-010994-8 (Berlin) ISBN 0-89925-203-6 (New York) NE: Sedlacek, Blahoslav [Hrsg.]; Microsymposium on Macromolecules

Copyright © 1986 by Walter de Gruyter & Co., Berlin 30. All rights reserved, including those of translation into foreign languages. No part of this book may be reproduced in any form - by photoprint, microfilm or any other means nor transmitted nor translated into a machine language without written permission from the publisher. Printing: Gerike G m b H , Berlin. Binding: Lüderitz & Bauer Buchgewerbe GmbH, Berlin. - Printed in Germany.

PREFACE

In this volume, the reader will find the majority of short special lectures and papers based on posters contributed to the 28th Microsymposium "Polymer Composites" held in Prague on July 8-11, 1985. This meeting, organized as part of the 1985 program of Prague Meetings on Macromolecules (PMM, including Microsymposia, Discussion Conferences and Summer Schools), was sponsored by the International Union of Pure and Applied Chemistry (IUPAC Macromolecular Division), the Czechoslovak Academy of Sciences (CSAS) and the Czechoslovak Chemical Society. The Microsymposium was organized by the CSAS Institute of Macromolecular Chemistry (V. Kubanek, Director; P. Cefelln, IMC Scientific Secretary and PMM Chairman) in cooperation with the CSAS Institute of Theoretical and Applied Mechanics (J. Nemec, Director). The scientific program of the meeting was prepared by J. Kolarik, Microsymposium Chairman, and members of the Programm Committee. At this meeting, 18 short special lectures and 46 posters were presented; however, numerous papers were contributed by those not able to attend the meeting. (Ten main lectures were published separately in Pure and Applied Chemistry, Vol. 57, No. 11, 1985.) The contributions deal with various topics: structure and properties of polymer matrix and interphase boundary; research methods; prospects of polymer composites; polymer matrix-polymer fibers composites; theoretical and practical scope of short-fiber application; molecular composites; fracture mechanics of composed polymer systems; hybrid composites and matrices; diffusion of lowmolecular-weight composites; particulate composites of polypropylene and polyethylene; etc. To locate a desired topic, please consult the Subject Index or the Contents which are divided into individual sections: 1. Introductory articles. 2. Preparation-properties-use. 3. Mechanical stability of composites. 4. Structures and methods. Papers of a rather general nature are included in the first section, articles of a specialized nature in the respective section.

VI

In spite of all our efforts, a number of mistakes and 1diviations' doubtlessly still remain in the text. Please excuse our faults (errare humanum est) - but not those of polymer composites. I would like to express my warm thanks to all the authors who contributed to this volume, devoting much time, energy and patience and, also, to de Gruyter Publishers and their leading coworkers for their unwavering interest and helpful approach in all complicated matters. Prague, April 1986

Blahoslav Sedlafiek PMM Editor

CONTENTS

INTRODUCTORY

ARTICLES

Molecular Composites M. Takayanagi

3

Modeling of Molecular and Particulate Composites T.S. Chow

19

Microstructural Considerations in the Development of Improved Fibre Reinforced Thermoplastics M.J. Folkes, S.T. Hardwick, and W.K. Wong

33

Mechanics of Hybrid Composites H. Fukuda

51

Conducting Polymer Composites A.T. Ponomarenko, V.G. Shevchenko, I.A. Tchmutin, A.A. Ovchinnikov, and N.S. Enikolopyan

67

Rheological Properties of Amorphous Polymers Containing Platelet Fillers R.P. Char toff and E.H. Eriksen

89

Comparison between Theoretical and Practical Mechanical Properties of Polyolefins-Glass Beads Composites M. Pegoraro

105

Oxidative Degradation of Polypropylene Catalysed by Mineral Fillers K. Veseltf, J. Petrflj, and A. Zahradnifikov^

123

PREPARATION-PROPERTIES-USE

1. Polyolefin Composites Some Properties of Highly Filled Oriented Polyolefins with Chalk and Other Systems M. Kryszewski, A. Gai^ski, and T. Kowalewski

141

VIII Cellulosic Fillers for Thermoplastics 153

C. Klason and J. Kubit The Effect of Fillers on the Rheological and Mechanical Properties of Polypropylene Composites B. Pukanszky, F. Tiidfls, and T. Kelen

167

Relation between Critical Fiber Length and Tensile Strength for Glass Fiber - Polypropylene Resin Composites M. Miwa and T. Ohsawa

183

Polypropylene-Mica Composites

J.P. Trotignon, J. Verdu, R. de Boissard, 191

and A. de Vallois Influence of Fillers on the

Process of Polypropylene

E.J. Paakkonen, S.N. Magonov, and P. Tormala

199

Polypropylene Composites - Dependence of the Yield Stress on the Concentration of Particulate CaCO-j Filler J. Hugo and M. Houskovi

207

Thermoelastic Effect of "Polypropylene - CaCO^" Composites. The Influence of the Composition, Rate of Strain and Temperature J. Hugo, M. Houskova, and V. Matena

217

Rigid Structural Foams from Composite Polypropylene/ Calcium Carbonate F. Smejkal

225

Instrumented Impact Studies of Some Thermoplastic Composites H. Hoffmann, W. Grellmann, and V. Zilvar

233

The Interactions on the Interface of Polypropylene and Organic Pigments A. Marcincin, E. Zemanovi, and J. Beniska

243

2. Other Polymer Composites Use of Wood Fibers in Thermoplastic Composites III. Polymethyl methacrylate B.V. Kokta, P.D. Kamdem, A.D. Beshay, and C. Daneault

... 251

IX

Structure and Properties of Polystyrene Prepared by Polymerization in the Presence of Carbon Fillers M.T. Bryk and A.F. Burban

269

Thermomechanical Behaviour of Graft Styrene Copolymers and their Composites M.C. Michailov and L.I. Minkova

275

Composites of Alkaline Poly(6-Caprolactam) Polymerized in Situ J. Horsky and J. Kolarik

283

Anionic Poly (6-Caprolactam) Composites Polymerized in Rotating Moulds J. Horsky Mechanical Properties of Soft PVC-Textile Composites J. Pigiowski and M. Koziowski

291 297

Hydrophilic and Thermoformable Silicone Rubber Composite P. Vondrâcek, J. Hrudka, J. Sulc, and P. Lopour

303

Effect of Polymer Matrix on the Efficiency of Microcapsulated Flame Retardants T.V. Popova, R.P. Stankevitch, M.S. Vilesova, N.A. Khalturinskii, and A.A. Berlin

311

Properties of Protein Modifications Covalently Linked to Particles J. Zemek, L. Kuniak, I. Novak, and D. Berek

323

3. Thermosetting Matrices The Reversibility of Hydrothermal Effects in FibreResin Composites G. Pritchard and S.D. Speake

329

Low Temperature Relaxation Behavior of Epoxy Resins T. Takahama, C.S. Wu, A. Chen, S. Pangrele, and P.H. Geil

347

Thermally Stimulated Depolarization of Radiation Cured Unsaturated Polyester Resin-Glass Micronodules Composites Z. Jelcic and F. Ranogajec

3 63

X

Glass Beads Filled Epoxy System: The Toughening Effect of an Introduced Elastomer Interphase Y.G. Lin, J.P. Pascault, and H. Sautereau

373

Structure and Mechanical Properties of Polymer-Phenolic Microsphere Composites D. ¿uchowska, J. Malczewski, and L. Wozniak

381

Catalytic Effect of the Solid Acids at Amino Resins Solidification V.M. Tcheshkov, M.M. Natova, G.Z. Zachariev, G.V. Kozlov, and T.M. Morozova

389

MECHANICAL STABILITY OF COMPOSITES 1. Interface Effects Interphase Effects on Viscoelastic Properties of Polymer Composites F.H.J. Maurer

39g

On the Influence of the Interface on Processing and Application Properties of Filled Polymers M. Ratzsch, H.-J. Jacobasch, and K.-H. Freitag

413

About Interface Problems in Kaolin-Filled Polyethylene J. Gahde

431

Structure-Property Relationship at Composite Interfaces J.D. Miller, H. Ishida, and F.H.J. Maurer

449

The Effects of Additives on the Structure and Mechanical Properties of Polypropylene-Filler Systems G. Marosi, G. Bertalan, I. Rusznak, P. Ana, and I. Molnar

457

The Properties of Polymers Containing Petroleum Stabilizers G.F. Bolshakov and A.A. Sidorenko

465

Filling of Polymers with the Aid of Coupling Agents K. Szi jar to and P. Kiss

473

XI

2. Failure and Fracture Effects Some Observations about the Failure of Polymer Composites J. Nemec

479

On the Effect of the Yarn Rupture Elongation Scatter on the Strength of Unidirectional Organic Fibre Reinforced Plastics S.L. Bazhenov, A.M. Kuperman, L.V. Puchkov, E.S. Zelenski, and A.A. Berlin

487

A Study of the Adhesion Strength in the MatrixPrestrained Fiber Joints A.M. Kuperman and Yu.A. Gorbatkina

497

Strain Rate Dependences of Deformation Behaviour and Fracture Surface Morphology in Polypropylene with Short Glass Fibres M. Sova

507

Fracture Toughness of Filled Thermoplasts at Dynamical Tests E. Nezbedova, J. Ponesicky, and M. Sova

515

STRUCTURES AND METHODS 1. Structure and Properties of Composites A Study of Reactive Polymer Morphology Using Holographic Microscopy R.P. Chartoff, H.G. Spicer, G. Kevin, and J.K. Johnson

525

Polyethylene - Polystyrene Gradient Polymer.II. Influence of the Diffusion of Styrene on the Structure of Host Polymer P. Milczarek and M. Krvszewski

531

Structure of Polypropylene/EPDM Elastomer/Calcium Carbonate Composites J. Kolarik and F. Lednicky

537

XII

Mechanical Properties of Polyethylene and Polypropylene Filled with Calcium Carbonate J. Kucera and J. Kolarik

545

Mechanical Properties of Three-Component Polypropylene Composites B. Pukânszky, J. Kolarik, and F. Lednicky

553

Structure of Polyethylene Filled with Mineral Filler and Glass Beads V. Svehlovâ

56X

2. Testing Methods Detection of Interfacial Debonding in ParticleReinforced Composites E.A.A. van Hartingsveldt

559

Three-Fibre Method for Measuring Glass Fibre to Thermoplastic Bond Strength P.A. Jarvela, P. Tôrmala, and P.K. Jarvela

575

Influence of the Viscoelastic Properties on the Bonding Strength of Metal-Polymer Composites A. Bauer and C. Bischof

5g3

Cafod - Computer-Aided Fiber Orientation Determination in Composites V. Djakovich, S. Fakirov, and L. Christov

589

Homogeneity of Polymer Composites F. Rybnikâr

597

Microscopic Methods Characterizing the Dispersion in Mineral-Filled Thermoplastics V. Svehlovâ

607

ABBREVIATIONS

615

AUTHOR INDEX

617

SUBJECT INDEX

619

INTRODUCTORY ARTICLES

MOLECULAR COMPOSITES

M. Takayanagi Kyushu Sangyo University, Faculty of Engineering Matsukadai, Fukuoka 813, Japan

Introduction Light weight, high strength materials are strongly desired especially in the field of aerospace engineering. Fiber reinforced composite materials are already contributing in this field. Recently, a new type of composite materials, being called "molecular composite" are gradually attracting the attentions in this field. Molecular composite is designed to use rigid rodlike molecules as reinforcement for the flexible molecules as matrix. The extension of the basic principle of fiber reinforcement to the molecular level has a possibility of providing new kinds of raw materials. Molecular composites have been explored by two groups: one is the author, using aramids, and the other is the group of U.S. Air Force Aeronautical Laboratory (1) and the staff of Dayton Research Institute (2). The latter group used wholly aromatic heterocyclic rod polymers such as poly(p-phenylene benzbisthiazole) (PBT) as reinforcement and heat-resistant resins such as poly(2,5(6)benzimidazole) (ABPBI) as matrix. They reported data on uniaxially oriented sheet of PBT/ABPBI = 30/70 that modulus E/strength a b is 120 GPa/1.4 GPa after annealed at 550°C (3). The data are comparable to those of aluminum of 70 GPa/0 .1 7 GPa with melting temperature of 660°C. Takayanagi et al. (4) used various conventional polymer matrices such as nylon, polyvinyl chloride, ABS resin, nitrile butadiene rubber (NBR) and poly(amide imide). It is noticeable that three-dimensional reinforcement is realized in their methods. In molecular composite (MC), the fineness of reinforcement is pursued to its limit, i.e. molecular dimension, otherwise microfibrillar dimension of 10 nm in diameter have been used. The rodlike molecules such as aramids and PBT retain their rigidity even as isolated molecules. Thus, these rodlike molecules dispersed

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

4 on molecular level are expected to reinforce the matrix flexible molecules if the molecular interaction between both components is strong enough. The merits of MC are expected in the large aspect ratio of rodlike molecules since the molecular size corresponds to that of reinforcement in an ideal case. Another possible merit is in the realization of ideal valence bond strength in the main chain . The macroscopic fiber is composed of bundles of fibrils (ca 100 nm in diameter), which in turn are composed of microfibrils (ca 10 nm in diameter). Many defects associated with the ends of fibrils and microfibrils reduce the fiber strength largely. Theoretical strength of polyethylene chain is 20 - 30 GPa, whereas the actual strength of PE fiber is 0.8GPa or less for normal fiber, and 3 - 4 GPa for ultradrawn PE fiber with ultrahigh molecular weight. The patent of MC was first applied by Takayanagi (9) in Japan, and the same idea, but using wholly aromatic heterocyclic molecules, was patent-applied by Helminiak et al. (1) in the United States, quite independently. Another reason for my starting the exploration of MC was found in superstructural characterization of crystalline polymers. Takayanagi et al. (5) proposed 1% fraction of tie link connecting the neighboring crystallites in fiber microfibrils, based on the analyses of viscoelastic relaxation curves of highly oriented and well annealed films, noticing their modulus anisotropy. Direct support to this view was given by Keith et al. (6). They cocrystallized polyethylene with n-paraffin which was removed by solvent after crystallization. Electron microscopic observation of the residue revealed that microfibrillar network of PE was developed in the texture, connecting neighboring spherulites and lamellae, while a major fraction of chains are folded and coiled to fill up the interspace of microfibrillar network. This structure is mechanically inefficient as external force is supported only by a very small fraction of extended chains. The appearance of Kevlar fibers by du Pont company in 19 73 was timely for our exploration of MC. Sulfuric acid dissolves poly(p-phenylene terephthalamide) (PPTA) and nylon 6 or 66 to form an isotropic solution below the critical concentration for liquid crystal formation. The ternary system was extruded in a coagulating bath and the MC of PPTA/nylon was first prepared by Takayanagi et al.

5 (7) . Sulfuric acid as a solvent decomposes conventional polymers used for matrix. A new and usable solvent for PPTA has been found and various MC's have been prepared by Takayanagi et al. (2) afterwards. The U.S. Air Force group prepared an isotropic solution of PBT and ABPBI in a mixture of methane sulfonic acid and chlorosulfonic acid, and the extrudate was coagulated in aqueous ammonium solution to obtain MC.

Preparation and Characterization of MC Coprecipitation method The isotropic solution of rodlike and flexible molecules in a common solvent was prepared at a concentration below the critical point for liquid crystal formation. Hwang et al. (2) showed that the phase diagram for the ternary system of solvent/PBT/ABPBI prepared experimentally well accorded with Flory's theory. The isotropic solution is extruded into a coagulant to avoid crystal formation as far as possible. If ideal system is realized, the rodlike molecules will be unimolecularly dispersed in the flexible moleculaes after coagulation. However, aramid molecules in MC of PPTA/nylon usually form microfibrillar network as mentioned later, the diameter of which was estimated to be ^30 nm with high resolution electron micrograph (4,7,10). Copolymers of aramid tend to be dispersed more homogeneously and the recent trend is along this line. The MC texture observed with a polarization microscope with crossed polarizers shows a wholly birefringent image caused by local orientation of flexible molecules associated with the rodlike molecules. Microfibril itself cannot be discriminated with such optical method. To find out new solvent dissolving both components of rigid and flexible molecules is of primary importance. Takayanagi et al.(8) found that sodium hydride (NaH) and dimethyl sulfoxide (DMSO) form sodium methyl sulfinyl carbanion with generation of hydrogen gas and this reagent diffuses into freshly polymerized PPTA to substitute the proton in amide group of PPTA with sodium ion. Metalated PPTA dissolves in DMSO to form a red colored homogeneous solution. Reaction scheme is represented as follows.

6 CH, / 3 0=S + NaH \ CH

CH / „ » 0=S +H \ - + CH Na

PPTA „ >

+

Na

N "-CO-Q-CO^ + Na

DMSO is a strong solvent for various flexible polymers and an isotropic solution of ternary system is easily prepared. Fig. 1 shows polarized optical micrographs of the MC of PPTA NBR (a) and chopped strands-reinforced NBR (b) with crossed polarizers .

100pm ,

100pm (a)

(b)

Fig. 1. Polarized optical micrographs with crossed polarizers of (a) vulcanized molecular composite of PPTA/NBR= 10/100, and (b) macroscopic fiber composite of Kevlar and NBR at the same ratio of PPTA/NBR as in (a) (Takayanagi et al. (10)).

NBR the nitrile butadiene rubber, with acrylonitrile content of 40%; NBR is dissolved in dimethylformamide (DMF). The solution of NBR in DMF is homogeneously mixed with metalated PPTA in DMSO to form an isotropic solution, which is coagulated with acidic water. The sample in Fig. 1(a) is a vulcanized MC of PPTA/NBR= 10/100 (6.4 vol % PPTA). Fig. 1 (b) is the macroscopic fiber composite with

7

the same composition as in (a), in which the vulcanized NBR matrix shows dark area due to its optically isotropic nature under crossed nicols. No fibrillar systems can be detected at this magnification for MC. It may be said that a new raw material has been born from NBR and PPTA. As easily inferred from the micrographs, the mechanical properties of the MC are quite different from those of macrofiber composite. Especially, crack growth is strongly impeded in MC and the tear strength is remarkably improved (10). Interfacial interaction is important in MC as in macroscopic composites. Accelerated crystallization of nylon 6 in presence of PPTA was proved by various facts: (1) the fractography reveals that PPTA microfibrils with 30 nm in diameter are broken in splitting, being protruded from fractured surface of the MC; (2) the DSC curve in cooling process, from molten state, shows the beginning of crystallization at the temperature close to the melting temperature of nylon 6; (3) the wide angle X-ray diffraction of the quenched MC from the melt indicates the appearance of the a-form crystal, while that of nylon 6 shows diffused diffraction; (4) the isothermal crystallization process measured by dilatometry shows that the incipient crystallization is found at the early stage, and Avrami's constant is 1 2, while that of nylon 6 is 4; and (5) the epitaxial growth is observed by SEM at the interfacial boundaries between Kevlar filament surface and nylon 6 (11). It was concluded that a very small difference in the distances between neighboring hydrogen bonding sheets in crystals of Kevlar and nylon 6 allows the epitaxial growth of nylon 6 on Kevlar surface. In the case of PPTA and polyvinyl chloride, it was proved by using benzanilide as a model compound of PPTA that the IR band characteristic of amide group hydrogen bond shifts in presence of benzanilide, due to the strong molecular interaction between both components. The method using a suspension of aramid fibrids A very thin layer of PPTA solution in sulfuric acid was smeared on a glass plate and soaked in aceton, being irradiated by ultrasonic wave. The fibrids with several tens nm in diamater are obtained as an aceton suspension. By replacing aceton with a tetrahydrofuran solution of polyvinyl chloride and solvent-casting, a composite film of PPTA fibrid and PVC can be prepared. When the en-

8 tanglements of fibrils are formed in the composite film, the SEM observation of fractured surface shows formation of large holes, reflecting heterogeneous dispersion of microfibrils of PPTA in PVC. In principle, it cannot be expected to prepare the ideal molecular composite with this method. Copolymerization methods Apart from the simple method of physical blend of straight aramids and flexible molecule matrices, preparation of the copolymers linking rigid rodlike segments to flexible chains according to the molecular design is a more hopeful method. Homogeneous dispersion and improved compatibility might be realized with this method. Several examples are cited : (1) Block copolymerization of nylon 6 or 66 with PPTA is performed with various methods (7). For example, triblock of poly(e-caprolactam-b-p-phenylene terephthalamide) is prepared by the following scheme:

H^N—(—nylon

6-)-C00H +

ClCO-^^-COCl

N

^

C1C0O-^^-CO—NH-nylon

>

6-)-C00H

I /v I + H2N-Q-NH2 + cicoT>coci

nylon 6-b-PPTA-b-nylon 6

(2) Polymerization of styrene in presence of I ^ N — S - S — N H ^ gives H 2 N—(-PSt-J—NH 2 (II). C1C0—f-PPTA-)—C0C1 (III) is synthesized with the conventional method, and the reaction of II with III gives a multiblock copolymer of —f-PPTA-b-PSt-}-^ . When H ? N— SH is used, a triblock of the same type is prepared. (3) Polybutadiene (PBD) with both chain ends being C0C1 in benzene solution is reacted with a solution of amino group-terminated PPTA in N-methyl pyrrolidone (14). Block copolymer of BDR-B—PPTA prepared with this method shows a remarkable toughening effect for ABS resin (15). (4) To the metalated PPTA solution in DMSO, a polyamic acid (PAA) solution in DMSO is added to obtain the MC of PPTA/ PAA, which is further cured at 260°C to convert PAA to polyamideimide (PAI) (12). Instead of straight PPTA, copolyamide of poly

9 (p-phenylene-co-3,4'-oxydiphenylene terephthalamide) was found more effective in increasing the ultimate elongation from 17% of MC of PPTA to 86% of MC of copolyamideimide (12).

500nm

500nm

Fig. 2.Electron micrographs of molecular composite of PPTA/nylon 6 =7/93. (a) Fracture surface. Triangular shadows are associated with each hole with 30 nm in diameter. (b) Nylon 6 matrix was removed by dissolving with formic acid and the microfibrillar network of PPTA remains (4,7).

Evaluation of tensile moduli of molecular composites In the electron micrographs of the fractured surface of the MC of PPTA/nylon 6, the aspect of fractured PPTA microfibrils can be observed (Fig. 2(a)). Diameter of microfibril is about 30 nm when high molecular weight of PPTA is employed. Fig. 2(b) shows the remaining microfibrillar network of PPTA after removing nylon 6 with formic acid from the MC. Microfibrils of PPTA were also confirmed in the MC of PPTA/NBR after extraction of NBR with DMF. The calculation of modulus of the MC as characterized above is somewhat different from that applied to the system reinforced with macroscopic fibers, which are dispersed separately without forming a network. The lattice model is assumed here as shown in Fig. 3 to comply with the actual state of the MC. Fig. 3(a) gives the upper bound of modulus as represented by eq.(1). V

E

f

V

f

c = ^ r E f + < 1 ~ - f >Em

(1

>

10 The model as shown in Fig. 3(b) gives the lower bound as represented by r

v

f

1

-

v

f 1

where E c , E^ and E m are the moduli of composite, fiber, and matrix, respectively, and V^ is the volume fraction of fiber. For random-oriented, three-dimensional fiber composites, the tensor analysis gives the coefficient of 1/6 instead of 1/3 in eq.(1) (13). With lattice model, the coefficient for two-dimensionally reinforced system is 1/2, and that of unidirectionally reinforced

(a) Upper Bound '-Vf, v,

irr

e 2 1 3 3

Ec=^-Ef+(l-f)Em

( b) Lower Bound

i-v f ^ f ^ l p lvt 2, 3

r= J Yi ' c I (lEi ++ § E m )

1 3

1-Vf tm

Fig. 3 . Calculation of modulus of molecular composite based on the lattice model (4). (a) Model providing the upper bound of modulus, and (b) the one for the lower bound.

11

system is 1, whereas with the tensor analysis for random-in-plane system, the coefficient is 1/3 and for unidirectional one the coefficient is 1. Informal data for the MC of PBT/ABPBI = 30/70 taken from Adams et al. (3) give the modulus of the unidirectionally oriented sheet being 120 GPa and that of the random-in-plane orientation being 62 GPa. The lattice model is more preferable to explain these data. However, more strict analysis is desirable for future discussion on this subject. To interpret the temperature dependence of viscoelastic relaxation curves of the actual systems, it is necessary to assume that the continuity of microfibril in each filament of lattice is interrupted by a small fraction of the matrix less than 1%, which results in reasonable explanation for modulus relaxation curves. It is conceivable that the microfibrils crystallized out from the matrix in MC will have modulus value lower than that of highly oriented and annealed macroscopic fiber. The lattice model is convenient to evaluate the modulus anisotropy generated during processing. For example, the milled and vulcanized MC of PPTA/ NBR (6.4 vol% of PPTA) shows the modulus anisotropy, which is more remarkable with increasing molecular weight of PPTA. The microfibrils tend to align along the milling direction. The higher the molecular weight of PPTA is, the stronger the shear stress along the milling direction during processing. The number of fibrils per unit cross-sectional area is increased along the orientation direction. Fig. 4 shows the relaxation curves for the MC of PPTA/NBR which is milled in an open roll and vulcanized with standard recipe in absence of carbon black. The molecular weight of PPTA are 4900 and 21900 for Figs. 4(a) and (b), respectively. The lattice constants of a, b and c are numerically evaluated for given diameter and volume fraction of PPTA microfibrils, referring to the close correlation of the modulus anisotropy and the swelling ratio anisotropy evaluated by soaking the sample in dichloromethane (10). Smallest area of lattice is the bc-plane, with the b-axis being the milling direction, and it means that the highest concentration of microfibrils is found along the a-axis direction. The swelling ratio is also the lowest along the same direction. The examples cited above belong to the two-phase system and their moduli cannot overcome the upper bound given by eq. (1). Moduli

12

My=

21900

E> 1000

A l l tul m

—

Ml

£ 2

Calcd Experimental

100

! \\ 10

NBR GUM STOCK

,1

1

240

1

260

1

320

1

360

N B R GUM STOCK

400

Temperature / K

2~0

280

(a)

320

Temperature/K

360

400

(b)

Fig. 4. The viscoelastic relaxation curves of milled and vulcanized MC of PPTA/NBR=6.4/93.6. Molecular weights of PPTA are 4900 for (a) and 21900 for (b). The lattice models employed for calculation of modulus anisotropy are indicated in the figures. The a-axis is the milling direction, the b-axis along the transverse direction and the c-axis along the thickness direction (10). of several systems employing copolymers with rigid chain sequence frequently overcome the upper bound.

In such cases, it is neces-

sary to take into account the effect of molecular interaction between both components.

Stress-strain behavior of molecular composites The initial slope of stress-strain curve, or Young's modulus is increased by incorporating rigid rodlike molecules in the matrix of flexible molecules, as far as the interaction at the interface of both components is strong.

Especially when the matrix is duc-

tile, the yield stress of MC is raised conspicuously by the small

13

STRAIN Fig. 5. Stress-strain curves for molecular composites for PPTA/ nylon 66=5/95 (curves A and B) and PPTA-b-nylon 66/nylon 66 with the same content of PPTA as in A and B (curves C and D). Curve E is for nylon 66 (7).

Fig. 6. Stress-strain curves for (A) molecular composite of simple blend of PPTA (2.5 wt%) and ABS; (B) molecular composite of PPTAb-PBD and ABS adjusted to 2.5wt%PPTA; and (C) ABS resin (15).

14

amount of rodlike molecules. In general, the ultimate elongation is remarkably decreased for the MC prepared by the coprecipitation method of straight aramid and matrix of flexible molecules (7). The curves A and B in Fig. 5 are for the MC prepared by the coprecipitation method with PPTA/nylon 66 =5/95, in which the molecular weights of PPTA are 4 500 and 980, respectively. Curves C and D are for the MC prepared by using the block copolymers of PPTA and nylon 66 with nylon block contents of 28 % and 37 % , respectively. The composition of the MC is adjusted to PPTA/nylon 66 = 5/95. The molecular weights cf PPTA in block copolymers are equal in both pairs of A and C, and B and D, respectively. The increased elongation of curves C and D in post-yield region is ascribed to the fact that the nylon blocks in copolymers are anchored into the nylon matrix. The removal of defects associated with the ends of PPTA microfibrils by block copolymerization has great influence on modulus, yield stress and toughness represented by strain energy at break, which are largely improved (7). Another interesting example using block copolymers is the MC of PPTA-b-PBD/ABS resin (14). ABS resin is biphasic system, in which the rubber phase of PBD is morphologically proved to be uniformly dispersed in the plastic phase of styrene acrylonitrile copolymer. The curve C in Fig. 6 is the stress-strain curve for the single system of ABS resin. By incorporating the PPTA-b-PBD in the ABS matrix to adjust the composition of PPTA to 2.5wt% with the coprecipitation method employing DMSO as a common solvent for ABS and metalated PPTA copolymer, the stress-strain curve B can be obtained. Strain energy to break as a measure of toughness is increased by about 3.7 times with respect to that of single ABS resin (curve C). Simple molecular composite of straight PPTA and ABS shows little increase in yield stress, but the ultimate elongation is largely decreased as expected from the effect of defects at the fibril ends. The SEM observation of the side surface of the film of the MC using block copolymer fully stretched close to the fracture reveals that the craze generation is uniformly spread all over the surface. PBD block in copolymer has an affinity to PBD phase in the texture of ABS and the PPTA block in copolymer has an affinity to the plastic phase in ABS resin at the same time. Rigid rodlike PPTA segment dispersed uniformly

15

in the texture of ABS resin will be effective in impeding the crack propagation. The third example is the reinforcement of PPTA sequence for polyamideimide (PAI). Both components of PPTA and polyamic acid (PAA) as a precursor of PAI are dissolved into DMSO using metalation reaction. Stress-strain curve of the MC of PPTA/PAI show stiffer and stronger but less elongation than that of PAI itself. Thus, the copolymerization method has been applied to improve the ultimate elongation. Curing of PAA to convert to PAI at 120°C and finally 260°C is represented by +£>-£>™-CO-£>CO-NH+n C0-NHo I 2 C

.curing

NH-CO o c

N

N-

As a molecular reinforcement, the following random copolymer (III)

Fig. 7. Stress-strain curves of polyamideimide(PAI) (curve 1), and the molecular composite of PPOT/PAI=30/70 (curve 2) (12).

16

is employed to improve the dispersion of rodlike PPTA segment in PAA :

III PPOT (x/y = 50/50) Polyamic acid I is dissolved into N-methylpyrrolidone (NMP). Copolyamide III, poly(p-phenylene-co-3,41-oxydiphenylene terephthalamide) (PPOT) is dissolved into NMP added with CaCl 2 . Both solutions are mixed to form an isotropic solution and solvent-cast to form a film of I and III. By curing the composite film at 120°C, 150°C and finally 250°C with gradual raising temperature, the MC film of II and III with the ratio of 111/11=30/70, is prepared. A remarkable feature of this MC is found in the temperature range around 300°C, in which region the modulus of MC overcome the upper bound modulus value given by eq.(1). This means that the strong molecular interaction is generated between reinforcing molecules and matrix flexible molecules. The stress-strain curve 2 in Fig. 7 shows that the ultimate elongation amounts to 86 % with yield stress being slightly higher than that of PAI film as a reference (curve 1). The ultimate elongation of PAI is 17 % . Hence, strain energy to break or toughness is increased almost five times. Modulus increased from 2.6 GPa to 3.6 GPa . The improvement of mechanical properties as found in this case can be realized by increasing the miscibility of both components with the copolymerization method.

Conclusion The MC shows remarkable improvement in modulus, ultimate strength, and heat resistance when suitable conditions are satisfied. To achieve such a success, the reinforcing molecule should be rigid rodlike, being geometrically linear. Interfacial adhesion is a governing factor for effective reinforcement. Molecularly uniform dispersion of rodlike molecules is most desirable due to the increase in aspect ratio and the removal of vital defects. Simple blends of rodlike and flexible molecules tend to form a microfibrillar network of rodlike molecules in the texture of MC.

17

Modulus calculation of such a system is amenable- with quasi-threedimensional lattice model, which is applicable even to the system with modulus anisotropy. A new solvent of aramid using sodium hydride and DMSO was found, which enabled us to prepare various kinds of MC's. Vital defects associated with chain ends can be removed by the copolymerization method, which is helpful in increasing the miscibility of reinforcing rodlike segments. As a result, the ultimate elongation or toughness is largely improved. The modulus for the molecularly mixed system overrules the upper bound value for the two-phase system. The remaining problem for the future is to modify the matrix in addition to the modification of rodlike sequence. The processability of MC's is another subject to be solved. To employ a more fluid curable resin for the matrix is a possible way. The U.S. Air Force group (3) succeeded in realization of high modulus and strength comparable to those of aluminum. However, the ultimate elongation is only 1.5 % . The sheet structure has already been set. If easy processability as found in aluminum can be provided to the MC, then it may be said that plastics replacing metal are actually realized.

References 1. Helminiak, T.E. 1978. U.S. Pat. application:902 525. 2. Hwang, W.-F., D.R. Wiff,C.L. Benner, T.E. Helminiak. 1983. J. Macromol. Sci.-Phys. B 1_7, 591. 3. Adams, W.W., T.E. Helminiak, A. Visvanathan. 1982. Proceedings IUPAC Macro'82, p. 828. 4. Takayanagi, M. 1983. Pure & Appl. Chem. 5j>, 81. 5. Takayanagi, M., K. Imada, T. Kajiyama. 1966. J. Polym. Sci. C (15), 263. 6. Keith, H.D., F.J. Padden, Jr., R.G. Vadimsky. 19 66. J. Polym. Sci. A-2, 4, 26 7. 7. Takayanagi, M., T. Ogata, M. Morikawa, T. Kai. 1980. J. Macromol. Sci.-Phys. B 1_7' 5 9 1 • 8. Takayanagi, M. , T. Katayose. 1981. J. Polym. Sci.: Chem Ed., 19, 1133.

18

9. Takayanagi, M. 1977. Japan Pat. application S.52-131 436 ; U.S. Pat. application 958 324 . 10. Takayanagi, M., K. Goto. 1985. In: Molecular Characterization of Composite Interfaces (H. Ishida, G. Kumar, eds.). Plenum Press, New York-London, p.247. 11. Kumamaru, F., T. Oono, T. Kajiyama, M. Takayanagi. 1983. Polym. Comp. 4, 135, 141. 12. Mitsutake, T., K. Yamada, M. Takayanagi. 1984. Polym. Prepr. Polym. Soc. Japan 3_3» 796. 13. Garg, S.K., V. Svalbonas, G.A. Gurtman. 1973. In: Analysis of Structural Composite Materials. Marcel Dekker, p. 22. 14. Takayanagi, M., K. Goto. 1984. J. Appi. Polym. Sci., 2j), 2057. 15. Takayanagi, M., K. Goto. 1984. J. Appi. Polym. Sci., 2£, 2547.

MODELING O F MOLECULAR AND PARTICULATE COMPOSITES

T.S. Chow Xerox Webster Research Center, 800 Phillips Road 0114-21D Webster, New York 14580

Introduction A unified understanding of the structure-property relations for molecular and particulate comppsites is presented on the basis of our composite model (1) and kinetic theory of the glass transition (2). The composite model is extended to interpret the change of phase domain in polymer blends and the variation of internal stress versus crystal shape in semicrystalline polymers.

The effects of quenching and annealing on the mechanical

properties of polymers and composites are treated as a nonequilibrium relaxation phenomena.

structural

We then analyze the effects of fillers and plasticizers on the

relaxation in composites.

Particulate Composites We have analyzed the effects of filler composition, size and shape and interfacial adhesion on the tensile strength and modulus of two phase anisotropic heterogeneous materials. The particle shape effect is characterized by the ratio of major to minor axes p = c / a of a spheroid. For a uniaxially oriented structure, the tensile modulus E|| is given by ( E | | - E m ) / E m 0 = [ E|| ] = " < e T 3 3 >/e 3 3 A

(1)

where E m and E r refer to Young's moduli of the matrix and filler, 0 is the volume fraction of filler and e 3 3 A is the uniform strain applied to the system.

< e 3 3 T > is the volume

average of transformation strain which can be obtained by solving the generalized Eshelby equations (1). A typical representation of Eq. (1) is given in Figure 1. A comparison of the longitudinal Young's modulus of a material filled with fibers aligned in the direction of the applied tensile load is given in Fig. 2 for the Cox (3), Halpin-Tsai (4) and above equations. The materials are two filled and crystalline polymers that have a

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

111

10

§

:

10

i

i

r ~ r 1 M 11

r^su-

Ej/Em ASYMTOTIC VALUES

-

X V

r i Mil.

1

En/Em

— — -

NN

> -I £

-

^

/

-

0= 0 . 6 0 — y / - p*

-

0*0.20— "

i

"

i

—

—

—

i i i i > 11

10'

I

i

, i i i i i il

.

,

1 11 1 11

10 ASPECT RATIO p = C/o

Figure 1.

10

Relative Y o u n g ' s M o d u l i .

Ef/Em=2l.2 = 0 . 4 0

— — -I

I

L_l 10

Figure 2.

COX CHOW HALPIN-TSAI

I I II I 10

ASPECT RATIO

'

I I III' 10"

Comparison o f the C o x , C h o w a n d Halpin-Tsai equations.

21 ratio E f / E m ranging from 21.2 (glass in epoxy resin), 100 (boron in epoxy resin) to 2 4 0 0 (semi-crystalline polyethylene) at constant 0 = 0 . 4 0 . little difference among the three relations. increasing values o f E f / E m > 10 . 2

F o r lower values o f E f / E m , there is

T h e discrepancy among them increases with

Porter and his coworkers ( 5 ) have found that the

calculated aspect ratio o f 100 for crystallites in utra-oriented semicrystalline polyethylene from the Halpin-Tsai equation is inconsistent with observed value.

A calculation

of

crystalline shape by Eq. (1) gives p — 25 which compares well with the observed ratio o f 20 for the crystals o f the core region.

Figure 2 also reveals that the shape dependence o f the

tensile modulus starts to diminish for p > E f / E , n . Usually the primary failure mode o f a filled polymer composite involves

interfacial

debonding as discussed by Trachte and DiBenedetto (6). Based upon Griffith's concept o f fracture, a necessary, but not necessarily sufficient, condition for circumferential crack formation is that the strain energy stored in the two-phase medium must b e equal to or greater than the energy required to create new surfaces; Le., Aw =

- tt

f

e; j T ( r ) d v >

y

A

(2)

V

where r -

rf +

Ym—?fm's

polymer-filler interfacial energy per unit area (work o f

adhesion) and A is the surface area produced. formation

o f an

equilibrium

crack

is due

T h e strain energy associated with the to the

inhomogeneity and the uniformly applied stress

=

(8)

R

To solve eqs.(6), the following expressions are introduced:

u

™(5)

u

*> (5)

=

c

+

L

k

*

°° =

c +

k

L

+

v ° >

k

*

L

V k < ? '

v ° >

°°

Vk'f)

u

k(0)

+

(9) Vk

v°>

where V, V*, W, and W* are the influence functions.

Considering

the second conditions of eqs.(8), eqs.(9) become V ? )

=5

+ 1 v m _ k (?) u k (0)

+

I*

wm_k(a

D *(0)

(10)

=

? +

where g and g

I Vk'«>

V

0 )

+

I*

V

0 )

indicate, respectively, the summations with respect

to discontinuous LE and HE fibers.

Then the first conditions of

eqs.(8) lead to dv

1 + I k

m-k

(0)

d?

* dW m . (0) * UK (0) + I u*(0) = 0 K k d5 (m = broken LE fiber) (11)

1 + I k

dV* , (0) d?

uK (0)

4

+

f k

m

,

-k d5

t

• U K> )

= 0

(m = broken HE fiber) These equations are used later to determine

and U^CO).

56 Finally, the following two sets of the influence functions are obtained: (I)

d2v

m =

2V + V + V , = 0 m m m-1

d 2

*

V m R o ^^

(12)

* * 2V + V + V = 0 m m+l m

with the conditions of V in(0) = 1

(m=0),

d

> >

d v

V " >

dÇ (II)

V m (0) = 0 (m^O),

V*(0) = 0 m

u

dÇ

d2W * * ¿i - 2Wm + W m + W m-1, = 0 2

R

(13)

„

( u )

*

d W * 7" " 2W + W . , + VJ = 0 m m+l m dÇ

with the conditions of W m (0) = m0,

W*(0) = 1 (m=0),

W*(0) = 0 (m^O) m

*

dW (

C 3 D = - C 4 D = (2-A2) (2-A2) K

K = 1/(A 2 - A 2 )

Substituting eqs.(21) into eqs.(17) and considering a, b, Xi, , Ci , C2, and CaD being even functions with respect to 8, the following results are obtained. 1 r 7T - / tC 1 exp(-A 1 ?) + C 2 exp(-A 2 C)]cosme de

-

k

(23)

/„

C

3D

" ex P ]

^llZoU^ ^ de

The SCF of the m'th group of fibers is defined as P m (0)/P(«°) * * m m or P m (0)/P*(®). These are: m m p 0 m< ' _J? =1 P (00) m

+ 1 k

dV

, (0) d?

u. (0) + f k

* dW m , (0) * —El* U*(0) d? (24)

58 dV p (oo)

1 + I k

, (0) UK (0) + I k

d?

*

m-K d?

V

0 )

The present theory is an advancement over Hedgepeth's theory. Thus it should include Hedgepeth's theory as a special case of R=l, meaning one type of fiber. Suppose, for example, only one LE fiber of m=0 is broken. In this case the fiber nearest to the broken fiber is an HE fiber of m=0. Substituting R=1 into eqs.(22) and (23), we get, after some calculation, dV Q (0) d?

dV*(0) d£

_4_ 3it

(25)

Since eqs.(ll) reduce to 1 + dVo(0)/dS-Uo(0) =0 in the present case, Uo(0) becomes jt/4. Finally, the SCF of a fiber nearest to a broken fiber becomes V°>

, . dV*(0) -0 = 1 + U Q (0) = 4/3 d?

(26)

which is identical to eqs.(2) for r=l. Although this is just one example, we can show that our analysis coincides with Hedgepeth's in the special case of R=l. When R^l, eqs.(23) are difficult to solve analytically and therefore, numerical integrations have been carried out. Figure 4 is the case of only one LE fiber being broken. In addition to the nearest fiber(HE), the SCF of the second-nearest fiber(LE) is also shown in the figure. In a glass/carbon hybrid, R will be l/2~l/3. Since the failure strain of carbon fiber is smaller than 5 3r that of glass fiber, carbon fibers will break at an early stage of loading. According to Fig. 4, the SCF of LE fiber is always smaller than 4/3. W 1 b I Therefore we can say that glass fibers have a role of a 1 ' 0 -02 • 05 -I -2 crack arrester in a -5 1 2 5 10 20 s t i f f n e s s ratio R glass/carbon hybrid. A Fig.4. Stress concentration factor. similar tendency holds also

59 when several fibers are brokent5], although details are not shown here. Let's now consider the case of Fig.2(a).

Interply Hybrid.

Here,

the equilibrium equations become d2u EA — ^ „2 dx

+ £h _ d

2

m+1

m-1

2

, m

Gd . n

+

* _ m

m (27)

*

* fid m Gh * * E A + — (u + u , - 2um ) + n ( Um , 2 , m+1 m-1 dx d Although these *

*

d

u

* u j m

equations are more complicated than eqs.(4), we can obtain solutions in the same manner as used for the intraply hybrid sheet.

l.S

Figure 5 depicts the SCF in an interply hybrid in which one LE fiber is broken.

The

value D is d 2 / h 2

which 001

represents the spacing in the lamination

(see

_L 01

_L I R

_L 10

I

_L 100

Fig.S. SCF of interply hybrid.

Fig. 2). Intraply, Laminated Hybrid.

This is the case of Fig.2(c) and the

equilibrium equations are

EA

d2u £ + — ,2 , dx d ,2

E

* * A

d

U

Gd (u* + u* , - 2u ) + ^ m+1 m-1 m n

m

. Gh T < u m+l d

+

uj m (28)

*

dx

(u m

+

u

m-l "

2u

Bl

)

*

+

Gd IT

(U

m "

U

m}

* =

0

Figure 6 shows the comparison of the SCF between intraply and interply hybrid composites.

As far as LE fibers (B or D in Fig.6)

are concerned, an intraply hybrid is superior to an interply hybrid because the SCF of fiber D is always smaller than the SCF

60 of fiber B.

T

Statistical Calculation of Failure Initiation Knowing the above stress concentrations, failure initiation in an intraply hybrid

o-oi

sheetCFig.2b) is now predicted.

This is

an improvement to Zweben's

_L

Fig.6. Comparison of SCF between interply and intraply hybrid.

theoryt9]. The analytical model is shown in Fig.7.

The axial length of the

specimen is L and

is the so-called

ineffective length.

Thus each fiber consists

of M^ = L/6J j links.

B

The total number of

fibers in the composite is N, of which N/2

hA

are LE, and N/2 are HE fibers. We assume the cumulative distribution functions for the failure strains are

Fig.7. Model of failure initiation.

F (e) = 1 - exp(-pie ) F*(e) = 1 - exp(-rle s )

(29)

where p, q, r, and s are Weibull parameters and 1 is the length of link.

An asterisk(*) is again used for defining values related to

HE fibers.

It is natural to consider that an LE fiber will break

first, as was discussed previously. When the hybrid composite is subjected to a strain e, the expected number of scattered fiber breaks in the N/2 LE fibers is X

lh=I

M

h

N F




(30)

When an LE fiber is broken, a stress concentration will occur in the continuous fibers adjacent to the broken fiber, as is shown in Fig.4

*

Hereafter, we use the strain concentrations k, and kT h n instead of K L E and K h e which are identical under the linear elasticity condition.

61 The probability that the nearest LE fiber will break under the strain concentraion k^ is F(k. e)-F(e) l-F(e)

(31)

Since the model is symmetric, the probability that at least one of two nearest LE fibers will break is P

=

2h

1

( 1

-

"

V

2

(32)

and the expected number of these situations is X

2h - Xlh

P

2h

(33)

Zweben[9] defines the following condition to result in the failure of the composite: X2h(Ë2h)

=

1

(34)

where

the composite strain at which the nearest LE fiber

will break.

We follow this definition, although this criterion is

rather conservative.

This is similar to a di-plet[16].

Assuming 1-F(s)= 1 in eq.(31), P_, - 2P, in eq.(32), and a a exp(-pls q )=i-pi£ q in eq.(29), the failure strain is calculated to be e

=

2h

I

-

W (35)

In the case of a non-hybrid composite, e

2

=

[2»I.P2(^l)]-1/2q

is the lower bound (di-plet) of the failure strain. The hybrid effect refers to the condition in which the initial failure strain of a hybrid composite is greater than the failure strain of non-hybrid LE composite. observed by Hayashi[20].

This effect was first

Bunsell and Harris[21] attempted to

explain it from a thermal residual strain.

Zweben[9] and Manders

and Baderl221 also attempted to explain it from a statistical view point. In the present theory, the hybrid effect is

62 e,, = _£h

R £

e

=

[

6.

(k,q - 1) -l/2q h ]

(37)

2 6 (kq - 1)

2

where 6 and k are, respectively, the ineffective length and the strain concentration factor(=l.333) for non-hybrid composites. Although ZwebenC9] proposed a similar idea, he used kf instead of h k^ , by which the hybrid effect in a strict sense cannot be explained. Bunsell and HarrisC21] reports that the experimental hybrid effect for four-layer carbon/glass hybrid was R £ (exp) =1.31. Substituting necessary data into the present theory, R g becomes 1.11[71.

The present calculation does not account

for the effect caused by the residual thermal strain which is about 10% of failure strainC211.

If we add this to the above, the

hybrid effect would be modified to 1.21, which is fairly close to the experimental value.

Monte Carlo Simulation of the Failure Process It is not easy to analyze the

j

failure process of composites theoretically.

Hence, a Monte Carlo

simulation is sometimes used, as was mentioned in the Historical Review. Figure 8(b) shows a model for Monte Carlo simulation which is a type of Rosen model[111.

As the initial condition, a

stochastic strength, STR(i.j), is assigned to each link, (i,j).

A normal

distribution for the strength of links has been adopted; that is, STR(i,j) is a normal random number with a specified

(a) Fig.8. Model of Monte Carlo simulation.

mean value and standard deviation for the LE and HE fibers. A uniform strain is applied on the boundary of the composite model. 1.

The SCFs of all links, SCF(i.J), are initially assigned as

A link with the least value of STR(i,j)/SCF(i,j) is sought.

Let this link be (io.jo).

When this link breaks, stress

63 concentrations take place

0 0 0 0 0 0 0 0 0 0 0 0 0 3 0 0 0 0 0 0

in the 1 inks of i = io. Then SCF(io,jo) is replaced by a corresponding value which has been calculated

previously.

The smallest value of STR(i.j) is again sought

0 0 0 0 0 0 0 0 0 0 0 0 0 1 0 0 0 0 0 0

0 0 0 0 0 0 0 0 0 0 0 0 0 4 0 0 0 0 0 0

0 0 0 0 2 0 0 0 0 0 0 0 0 s 0 0 0 0 0 0

0 0 0 0 a 0 a 0 0 0 a 0 0 6 0 a 0 a c 0

0 0 0 a 0 0 0 0 0 2 a 0 0 0 0 a a 0 a 0 0 0 0 a a 0 0 0 0 0 10 0 13 0 14 a 0 il 0 0 4 0 0 3 0 7 IS > IS 1 0 a 0 0 0 0 0 0 9 a 0 0 0 0 0 0 0 0 0 0 0 a 0 0 0 0 0 a 0 0 0 0 a 0 a 0 a a 0 0 0 0 0 0 0 0 6 0 s 12

from the remaining MN-1 links.

LE HE LE HE LE

A link which has

the smallest value will break second.

hybrid

non-hybrid

Fig.9. Example of failure pattern.

This

procedure is repeated until all the links of the same i are broken. Numerical

calculations

have been conducted for a model of M=5 and N=20.

L E fiber c o m p o s i t e

t> è

hybrid

The average normalized strength and the standard

H E fiber composite

£ -5

deviation of both LE and HE fibers are assumed to be 1 and 0.1, respectively.

The ratio of

extensional rigidity, R, is assumed to be 1/3 which approximately

corresponds

to a glass/carbon hybrid.

_L normalized strain

E/E|_INK

Fig.10. Simulation of the stressstrain diagram.

Figure 9 is an example of the failure process. in a hybrid composite break at many locations.

The LE fibers

This is called a

multiple failureC211 which is common in a hybrid composite.

By

comparison, non-hybrid composites are apt to break drastically. Figure 10 shows an example of the stress-strain

relations.

This figure is very close to the experimental result of Bunsell and Harris[21],

It is reasonable that the ultimate failure strain

of a hybrid is larger than LE fiber composite and smaller than HE fiber composite. The rule of mixtures for hybrid composites can also be compared with the results of the Monte Carlo simulation.

The

64 solid lines of AC and CD in Fig.11 are the strength of hybrid composites predicted by the idea of the rule of mixtures.

On

the other hand, dots and vertical ranges show the average and standard deviation calculated by the Monte Carlo simulation.

For non-

hybrid composites such as GFRP and CFRP, the rule of

Fig.11. Rule of mixtures and Monte Carlo simulation.

mixtures predicts a slightly higher value than the value calculated statistically. The rule of mixtures should be modified by a factor k which is slightly smaller than unity.

In the case of hybrids, however, the

rule of mixtures seem to predict a lower strength than the actual strength.

Conclusions This paper first evaluates the stress concentration factors in hybrid composites. composites are treated.

Both interply and intraply hybrid Based upon the knowledge of the stress

redistribution at fiber breakage, two statistical approaches have been done.

One is to calculate the failure initiation.

modification of Zweben's theory has been conducted.

Some

Monte Carlo

simulation is another approach by which the failure mechanics of hybrid composites have been made clarified.

That is, the failure

of hybrids proceeds gradually due to multiple failure; its ultimate failure strain is higher than that of LE fiber composite; and the hybrid effect is explained quantitatively. Although the present analysis has been performed for a rather small model, the results offer a more comprehensive understanding of the mechanics of hybrid composites.

65

Acknowledgements The author wishes to express his sincere thanks to Professor Jack R. Vinson, University of Delaware, presently a Visiting Professor of the University of Tokyo, for his kind advice in preparing the manuscripts.

References CI] Fukuda, H. and T. W. Chou. 1981. Stress concentrations around a discontinuous fiber in a hybrid composite sheet. Trans. Japan Soc. Compos. Mater. 7, 37-42. [2] Fukuda, H. and T. W. Chou. 1982. A statistical approach to the strength of hybrid composites. Proc. ICCM-IV (T. Hayashi, K. Kawata and S. Umekawa, eds.). 1145-1152 [3] Fukuda, H. and T. W. Chou. 1982. Monte Carlo simulation of the strength of hybrid composites. J. Compos. Mater. 16, 371385. [41 Fukuda, H. and K. Kawata. 1983. A Monte Carlo simulation of the strength of laminated hybrid composites. Trans. Japan Soc. Aero. Space Sci. 25, 203-215. [5] Fukuda, H. and T. W. Chou. 1983. Stress concentration in a hybrid composite sheet. J. Appl. Mech. 50. 845-848. [6] Fukuda, H. and K. Kawata. 1984. Comparison of interply and intraply hybrid composites from the view point of stress concentration. Theo. Appl. Mech. 32, 459-466. [7] Fukuda, H. 1984. An advanced theory of the strength of hybrid composites. J. Mater. Sci. 19, 974-982. [81 Hedgepeth, J. M. 1961. Stress concentration in filamentary structures. NASA TN D-882. [91 Zweben, C. 1977. Tensile strength of hybrid composites. J. Mater. Sci. 12, 1325-1337. [10] Ji, X., G. C. Hsiao and T. W. Chou. 1981. A dynamic explanation of the hybrid effect. J. Compos. Mater. _1_5, 443461 . [11] Rosen, B. W. 1964. Tensile failure of fibrous composites. AIAA J. 2. 1985-1991. [12] Harlow, D. G. and S. L. Phoenix. 1978. The chain-of-bundle probability model for the strength of fibrous materials. J. Compos. Mater. 12, 195-214 & 314-334. [13] Zweben, C. 1968. Tensile failure of fiber composites. AIAA J. 6, 2325-2331. [14] Fukuda, H. and K. Kawata. 1976. Strength estimation of unidirectional composites. Trans. JSCM. 2, 59-62. [15] Bergman, B. 1981.

On the probability of failure in the

66 chaln-of-bundles model.

J. Compos. Mater. 15, 92-98.

[16] Batdorf, S. B. 1982. Tensile strength of Unidirectionally reinforced composites. J. Reinf. Plas. Compos. 1, 153-164 & 165-176. [17] Fukuda, H. and K. Kawata. 1977. On the strength distribution of unidirectional fiber composites. Fibre Sci. Tech. 10, 5363. [181 Oh, K. P. 1979. A Monte Carlo study of the strength of unidirectional fiber-reinforced composites. J. Compos. Mater. 13, 311-328. [19] Fukunaga, H., T. W. Chou and H. Fukuda. 1984. Strength of intermingled hybrid composites. J. Reinf. Plas. Compos. 3, 145-160. [20] Hayashi, T. 1972. Development of new material properties by hybrid composition. Fukugo Zairyo. 1, 18-20. [21] Bunsell, A. R. and B. Harris. 1974. Hybrid carbon and glass Composites. 5, 157-164. fibre composites. [22] Manders, P. W. and M. G. Bader. 1981. The strength of hybrid glass/carbon fibre composites. J. Mater. Sci. 16, 2233-2245 & 2246-2256.

CONDUCTING POLYMER COMPOSITES

A.T. Ponomarenko, V.G. Shevchenko, I.A. Tchmutin, A.A. Ovchinnikov, N.S. Enikolopyan Institute of Chemical Physics, Academy of Sciences, Moscow, USSR

Introduction The term polymeric composite materials (polymer composites) is kncwn to refer to heterogeneous systems that contain two or more phases, one of them being polymeric. Two or more-component structure of the composite is, strictly speaking, the basis of the theoretical possibility of developing such materials with a new set of properties. On the other hand, though the problem of replacement of the traditional materials by the composites is rather critical today, it would be naive to assume that the composites are developed only to replace other known materials. If one turns, for example, to the conducting polymer composites, it is necessary to emphasize that this field of organic material electronics furnishes the answer to a range of new technical problems; the analysis of literature shows that widening of the research can promote the development of materials with formerly unknown sets of properties.

1. Design of Conducting Composites 1.1 General principles and approaches Considering two or more-component structure of the composites, they are essentially hybrid systems. While "pure" commercial polymers refer to products that contain less than 5 weight % of additives, the composites, therefore, are materials that contain much greater amount of additives. At the same time, the composition, on the whole, determines the ultimate properties of the composite. Therefore, a possible starting point for conducting composites design can be variation of the amount of ingredients. The advantages of this method of combining the properties are the following (1): 1) extensive opportunities for the development of new materials;

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

68 2) many properties of the composites can be varied; 3) appreciable variations of properties can be found at small amounts of additives. However, this method has several disadvantages : 1) time-consuming and sometimes elaborate investigations are needed; 2) in some cases several properties can become better, while the others remain worse; 3) for some reasons multiple repetitions of the known solutions are possible. Another approach is possible: control of the ingredients and the composite structure. Positive factors in this case are the following: 1) many properties are varied; 2) transitions between the structures are possible at suitable influences. Disadvantages of this approach are the following: 1) the composite structure does not always determine its ultimate properties (e.g. "unsuccessful structure"); 2) adequate choice of the structure is necessary, incorporating its influence on the associated properties of the material; 3) formation of particulate structure is often complicated by the technological possibilities. Thus, we outlined general approaches to the composites design based on our own and literature experience (2,3). There is no reason to state that these conditions apply only to the design of conducting composites, they can be useful as a general guidance. In addition, it is necessary to emphasize that properties of the fillers and the media into which they are introduced, as well as distribution of the fillers in this media, are decisive for evaluation of electrical properties of the composite to be designed. Certain physical approaches to this problem already exist (3).

1.2. Major technological methods of conducting composites development In selection of the main course in conducting composites design the following factors must be taken into account: (a) potentialities of the composite and engineering problems that

69 are r e s o l v e d with

it;

(b)

technical

and e c o n o m i c a l

characteristics;

(c)

r e s e a r c h and d e v e l o p m e n t s p e n d i n g and, f u r t h e r ,

investments

in

production; (d)

i n t r o d u c t i o n of w a s t e l e s s

(e)

ecological

the

composite.

and s o c i a l

As t h e i n t r o d u c t i o n o f of

the composite,

cipal

physical

perties

of

spheres;

the a p p l i c a t i o n

can change a p p r o p r i a t e to consider

Table

1 presents

of

properties

the f o l l o w i n g

an a n a l o g y

f o r t h e problems o f

thermal c o n d u c t i v i t y ,

prin-

(4)

magnetostatics

and c a l c u l a t i n g

Electrothermal

its

pro-

i n terms

electrostatics,

of

elec-

and d i f f u s i o n ,

t h e p a r a m e t e r s t h a t should be o p e r a t e d i n c h o o s i n g

composite i n g r e d i e n t s 1.

in adjacent

parameters t h a t a l l o w t o determine the u l t i m a t e

conductivity

trodynamics,

Table

fillers

i s necessary

composites.

generalized indicating

it

technologies

consequences a f t e r

ultimate

the

properties.

Analogy

Electrostatics

Electrodynamics

Thermal conductivity

Efield

Eleetremotive force, E

•temperature gradient grad T

Field induction vector, D

Current density vector, j

Heat f l u x , q

Field induction vector B

Mass flux M

Dielectric permittivity e

Conductivity a

Thermal conductivity coefficient, A

Magnetic permittivity U

Diffusion coefficient

The s y s t e m a t i z a t i o n bridization

of

of

the i n g r e d i e n t s

properties

a multiparametric

is useful,

of

hy-

can be a c c o m p l i s h e d

in is

problem.

F u r t h e r we p r e s e n t p r i n c i p a l

perties

since

t h e optimum c o m p o s i t i o n

properties

of

fillers,

c o m p o s i t e t h a t a r e u s e f u l from b o t h r e s e a r c h and points of view.

Diffusion Partial pressure gradient grad P

Hfield

t h e above p a r a m e t e r s

many ways and so t h e d e t e r m i n a t i o n o f

Magnetostatics

These p r o p e r t i e s

the f i l l e r ;

filler

are the f o l l o w i n g :

packing;

m a t r i x and t h e

technological shape and

a d h e s i o n and o t h e r

pro-

inter-

a c t i o n s a t t h e i n t e r p h a s e boundary; e l e c t r o p h y s i c a l p a r a m e t e r s : -3 2 -1 -1 c a r r i e r and t r a p d e n s i t y , cm ; t h e i r m o b i l i t y , cm V s ; con-

70 tact types: ohmic, blocking, injecting, etc.; I-V characteristics; electrical conductivity a,ohm ^cm ^ ; temperature coefficient of resistance; percolation threshold; tensoresistance; dielectric constant e and loss factor tg6 ; damping, dB/m; photoconductivity; magnetoelectric properties; thermal electromotive force; gas ab- 1

- 1

sorption; thermal conductivity, W m K ; mechanical properties. It seems expedient to indicate the importance of complex investigations of both the ingredients and the composites for the determination of the above parameters, where the cooperation of chemists - physicists - technologists - consumers (in any order) is important. Thus, we have decided to design a new composite with a set of properties. What should we have at our disposal for this? Obviously, first of all the filler and the matrix with corresponding properties. The filler is characterized by the texture (dispersion, fibre, fabric, ribbon, special shape). It can be mineral, organic, metal, or carbon in various modifications. Organic matrices are the usual thermoplasts and thermosets while carbon, metal and inorganic matrices are also common. The filler and the matrix must be combined, i.e. a rational method for distributing the filler in the matrix must be chosen and utilized. This course in known to be extremely complex, especially in case of electronic materials. Now we briefly review the shape and geometrical parameters of fillers (5). Traditional shapes of particulate fillers are sphere, cube, parallelepiped or plate. These shapes are, naturally, idealized and the real particles are only approximated by them. Fibre fillers can be continuous, short or compound fibres. Fillers can be made of various materials: metals, carbon (quasi-metal), polymers (organic metals), composites (metal-coated polymer). Finally, the use of fabrics (carbon, metal or compound) for layered composites is known. Figure 1 presents typical filler packings. Particulate fillers can be used to develop matrix, statistical or structured composites. Version (3), for example, is interesting in that by selecting an appropriate filler or compounding conditions, favourable for the structure formation one can markedly decrease percolation threshold of the composite. Version (4) is also important, since the desired level of conductivity can be achieved at low conducting phase concentration by coating polymeric fillers with a conducting layer. The use of conducting fibres, fabrics or ribbons is helpful in developing layered or anisotropic composites with dif-

71

©© © © @ © © © © 1

Ill

II

6 7

Figure 1. Distribution of fillers in polymeric matrix 1 - matrix system, 2 - statistical distribution; 3,4 - structured composites; 5 - conducting fabric composite; 6,7,8 - conducting fibre composite; 9 - short fibre and particle composite; 10 - composite with particles of different size. ferent orientation directions. Combination of conducting fillers with different particle shapes, for example, fibre-sphere or two spheres leads to the appropriate hybridization of electric properties (Fig.1 ; 9,10) . Now we are coming to the technologies used to compound the ingredients in conducting composites production. They are: mixing of filler and matrix powders; compounding in the melt of thermoplastics; compounding in solution; compounding in the components of thermosets with subsequent mixing; impregnation of the conducting component with polymer and subsequent pressing; polymerization in the presence of fillers; electrochemical

polymerization on the filler

surface; deposition of radical, ionic or ion-coordination catalysts and subsequent polymerization of monomers on the filler

(polymeriz-

ation filling). Because of the variety of compounding techniques, the following important electrophysical properties are more or less accomplished: stable adhesion between the filler and the matrix; uniform proportion of filler and the polymer in composite and specimen; presence of necessary ingredients and inclusions: appropriate particle packing in the composite; longevity of the composites. So far we have treated general principles of composites design with specific electrophysical aspects in mind. Now we turn to particular electrophysical problems.

72

2. Charge Transfer in Polymeric Composites Because of their macroscopic dimensions (-10 ^m), conducting filler particles usually have properties similar to the bulk material. Composite becomes conducting when concentration of the filler exceeds a certain threshold value, which corresponds to the formation of an infinite chain of conducting particles in the heterogeneous material. Above the threshold the conductivity increases, approaching the conductivity of the filler due to the increase in the number of contacts between the particles. The value of the threshold concentration is determined by the statistical properties of the system and packing of the particles and thus depends on their shape and dimensions. Conduction mechanism below the threshold is more complex since it is governed by the injection of carriers into dielectric and their transport in it. In this region of concentrations conductivity of composites is varied widely from typical insulators to semiconductors. Extrinsic character of conductivity leads to non-linear current-voltage characteristics and the parameters of non-linearity, as well as conductivity values can be selected experimentally, thus making such composites interesting for dielectric electronics (6). The main problem in the synthesis of composites with conductivity below the percolation threshold is to ensure constancy and stability of electrical properties, which are hindered by the statistical variations in the thickness of insulating layers, presence of impurity levels and blocking barriers at the interphase boundaries. In general, conducting composites are a combination of randomly distributed conducting filler particles with dimensions far exceeding the mean free path of electrons and separated by insulating layers of matrix. Despite the different possible conduction mechanism, such disordered systems as composites can be treated in terms of general probabilistic and statistical approaches used in the analysis of transfer processes in disordered semiconductors and amorphous metals (7,8). The most general approach to the analysis of conduction in disordered systems is percolation theory (9-11). Two possibilities exist for flux to flow through a disordered system - diffusion and percolation, each of them being a statistical process. The difference between them can be readily seen on the example of random jumps in one-dimensional lattice.

73 In case of diffusion there is a non-zero probability of jumping to the right or to the left site in every point. In case of percolation certain sites can exist for which the probability of jumping to the neighbouring sites is exactly zero. The principal problem of percolation theory is to find the values C t (b) or C t (s), i.e. the minimum value of concentration C when as infinite cluster of conducting bonds (b) or sites (s) is observed. Thus, percolation threshold is the point at which a continuous macroscopic conducting chain first appears. For all the lattices studied percolation threshold for the bond problem is less than the value of threshold in the site problem: C t (s) 2C t (b). The values of C t differ even for the lattices of the same dimension. The difference is caused by the variation in the number of nearest neighbours z. Average number of bonds per site at the threshold zCt(b) is independent of the lattice type and depends only on the space dimension, for the two-dimensional case being zCt(b) = 2.0, while for three dimensions zC^(b) =1.5. For the site problem the difference of C t (s) values for various lattices is attributed to the fill factor f - the ratio of the volume of all spheres (with a radius equal to half of the distance between the nearest neighbours) to the total volume. If such spheres belong only to the percolating sites, the fraction of volume occupied by these spheres is equal to fC. The critical value of C t (s) is calculated for two-dimensional lattices from the condition fCt(s) = = 0.45 (1) and for three dimensions fCt(s) = 0.15 (2). The continuum percolation problem, which is most adequate to the conductivity of composites with random distribution of the filler particles consists in finding the critical fraction of volume occupied by the conducting regions, corresponding to the onset of current. For the approximate evaluation of C^ value the invariant of the lattice site problem fCt(s) ((1) and (2)) can be used, since fCt(s) is the fraction of volume filled by the spheres centered in non-connected sites of the lattice which corresponds to the onset of percolation through the contacting spheres (7). Percolation threshold for two dimensions is 0.45 and for three dimensions Cfc = 0.145. In two-component continuous medium that consists of filler particles with conductivity o^ and matrix particles with conductivity a- the distribution of particles is binomial (12):

74 P(a^) = C/f

and

P(a 2 ) = 1 - C/f

(3)

where C is the volume fraction of the material with conductivity a^ and P(a) is the probability of finding a particle with conductivity a at the given point. Depending on C, the conductivity of the medium varies from 0 2 at C < C^ when the filler particles are isolated to a^ at C > C ^ when an infinite cluster appears. If the variation of

a is relatively small, i.e. 1I (a„ - a )/a 1 | « 1 ' 1 m'' m

(a m

is

the conductivity of the media) which is true in the neighbourhood of the percolation threshold, the effective conductivity o^ can be found by perturbation theory

(13). In the vicinity of the threshold

a^ of a two-component mixture described by Eq.(3) equals: a

m

£

°m

s

am =

a

1(C "

C

t)B ' a

o2(Ct-C)~ 0lx«

,

,

C > C

t

C C^. These results indicate the existence of direct contacts between graphite particles. References 1. Stofa, J. 1979. Electrotechnical Materials - Questions and Answers (in Slovak), Bratislava, Alfa. 2. Shevchenko, V.G., A.T. Ponomarenko. 1983. Usp. Khimii 52!, 1336. 3. Philippov, P.G., V.G. Shevchenko, A.T. Ponomarenko, V.A. V. A. Bendersky, A.A. Ovchlnnikov. 1984. Electrical Properties of Composites with Conducting Disperse and Fibre Fillers (in Russian). N°1 (219). NIITEKHIM, Moscow.

86 4.

Chudnovsky, A.F. 1962. Thermophysical Properties of Disperse Materials (in Russian). Fizmatgiz, Moscow.

5.

Handbook of Fillers and Reinforcements for Plastics (H.S. Katz and V. Milewsky, eds.). 1 978. Van Nostrand Reinhold, New York.

6.

Ovchinnikov, A.A. 1983. Vestn. AN SSSR. N'1, 71.

7.

Shklovski, B.I., A.L. Efros. 1975. Usp.Fiz.Nauk, 117-, N'3, 401.

8.

Bonch-Bruevich, V.A. et al. 1981. Electronic Theory of Disordered Semiconductors (in Russian), Nauka, Moscow.

9.

Broadbent, S.R., J.M. Hammersley. 1957. Proc.Cambr.Soc. 53, 62a

10. Shante, V.K.S., S. Kirkpatrick. 1971. Adv.Phys. 20^, 325. 11. Kirkpatrick, S. 1973. Rev.Mod.Phys. _45, 574. 12. Shezman, A.D., L.M. Middleman, S.M. Jacobs. 1983. Polym.Eng. Sei. 23, 36. 13. Landau, L.D., E.M. Livshitz. 1959. Electrodynamics of Continuous Media (in Russian). Fizmatgiz, Moscow, p.67. 14. Electrical Transport and Optical Properties of Inhomogeneous Media (J.C. Gurland and D.B. Tanner, eds.). 1978. AIP, New York. 15. Bruggeman, D.A.G. 1935. Ann.Phys. (Leipzig), 24,

636.

16. Landauer R. 1952. J.Appl.Phys. 2^, 119. 17. Wiener, O. 1962. Abh. Sachs.Gesellsch. 32^ 509. 18. Yamaki, J., 0. Maeda, Y. Katayama. 1978. Rev.Elec.Commun.Lab. 26, 616. 19. Sheng, P. 1980. Phys.Rev. 21B, 2180 20. Truhan, E.M. 1962. Fiz.Tv.Tela _4, 3496. 21. Mellikhov E.Z. 1965. Fiz.Tv.Tela 7, 1529. 22. von Hippel, A.R. 1954. Dielectrics and Waves. Wiley, New York. 23. Bube, R. 1960. Photoconductivity of Solids. Wiley, New York. 24. Abeles, B., P. Sheng, M.D. Coutts, Y. Arie. 1975. Adv.Phys. 24, N* 3, 407. 25. Enikolopov, N.S. 1980. Priroda N"8, 62. 26. Aivaziyan, F.N., P.E. Matkovsky, A.T. Ponomarenko, V.G. Pavliv, I.F. Shamsullin, Yu.N. Kolesnikov, N.S. Enikolopyan. 1983. In: Complex Metaloorganic Catalysts for Olefin Polymerization (in Russian). Chernogolovka, p.73.

87

27. Galashina N.M., V.G. Shevchenko, P.M. Nedorezova, P.G. Philippo\$ A.T. Ponomarenko, L.N. Grigorov, V.l. Tsvetkova, V.A. Bendersky, F.S. Djachkovsky, N.S. Enikolopyan. 1983. In: Abstracts of the IUPAC MACRO-83. Bucharest, Romania. Sect. VI, p.121. 28. Galashina, N.M., V.G. Shevchenko, P.G. Philippov, A.T. Ponomarenko, F.S. Djachkovsky, V.A. Bendersky, N.S. Enikolopyan. 1983. In; Abstracts of VI International Microsymposium on Polymer Composites. Budapest, p. 39. 29. Sichel, E.K., J.E. Gittleman, P. Sheng. 1978. Fhys.Rev. 18B, 5712. ' 30. Resistors (in Russian) (Chetvertkov I.I., ed.). 1981. Energoizdat, Moscow. 31. Chung, K.T., A. Sabo, A.P. Pica. 1982. J.Appl.Phys. 53, 6867.

RHEOLOGICAL PROPERTIES OF AMORPHOUS POLYMERS CONTAINING FILLERS

R i c h a r d P.

PLATELET

Chartoff

The C e n t e r for B a s i c a n d A p p l i e d P o l y m e r R e s e a r c h , U n i v e r s i t y Dayton, Dayton, Ohio 4 5 4 6 9 , USA

Erik H.

of

Eriksen

Norsk Hydro, A.S., Porsgrunn,

Norway

Introduction F i l l e r s can bring a b o u t d r a m a t i c c h a n g e s in the l i n e a r

viscoelas-

tic p r o p e r t i e s of a m o r p h o u s p o l y m e r s a s m e a s u r e d by d y n a m i c mechanical methods.

The d y n a m i c m e c h a n i c a l p r o p e r t i e s

(including

the real part of the e l a s t i c m o d u l u s E', the i m a g i n a r y p a r t of e l a s t i c m o d u l u s E", a n d their ratio tan 6

= E " / E ' ) of a

u n c r o s s l i n k e d a m o r p h o u s p o l y m e r are i l l u s t r a t e d in Figure

1.

f r e q u e n c i e s of a r o u n d 1 to 100 Hz the p e a k s or d i s p e r s i o n s a n d t a n 6 o c c u r a t t e m p e r a t u r e s a s s o c i a t e d w i t h the g l a s s tion t e m p e r a t u r e T^ of the p o l y m e r .

the

typical At

in E"

transi-

The e n e r g y d i s s i p a t e d

into

heat for a u n i t d e f o r m a t i o n of a p o l y m e r is g r e a t e s t near T^ a n d a t l o w e r t e m p e r a t u r e s , since this h e a t g e n e r a t e d is to E".

Damping

g r e a t e s t near T

proportional

is m o s t r a p i d w h e r e the ratio E"/E' or

t a n

_l Z5 Q |io 6 t 2 70 - 72%; Na20 13-15%; Ca0 7-ll%; MgO 3 - 5%). Hydrolyzed glass beads: the beads were heated at 150°C in autoclave with water for 24 hours (6). Adhesion promoters: we used known ccmpatibilizers reacting with the glass on the surface, as stearylisocyanate (SIC) or isotropyltriisostearyl titanate (ITS) (4), and new ccmpatibilizers prepared by us, ref.5) as a random copolymer of acrylic acid and hexadecyl methacrylate or grafted polymers with controlled length of the grafted chains obtained by copolymerizing a macromer of 2.4-vinylpyridine with butylacrylate (7) in different configurations. All the adhesion promoters we used are endowed with reacting groups or polar groups having affinity for the glass, and methylenic groups having affinity for the matrix. The adhesive polymers were spread on the beads in solution: after evaporation of the solvent the beads were blended with pov*3ered PE. Samples for mechanical tests were 2 obtained by compression molding with a pressure of 50 kg/cm applied for 5 minutes at 140-160°C. Elastic moduli were measured at an elongation rate of 1 mm/min at 23°C and 50% r.h., using an Instron testing machine 1195 and extensometers having amplification ratio ranging frcm 100:1 to 1000:1. Other tensile tests were done with an elongation rate of 0.5 cm/min. Contact angles were measured with a Rame Hart goniometer, surface tensions with the same apparatus and the pendant drop method (8). Elastic moduli Stiffness is the most important property of engineering materials for mechanical design. In the cass of polymeric composites, linear viscoelasticity is always assumed to be valid for small deformations. Most of the studies are related to purely elastic composites. In principle, the solution of purely elastic problems

107 can be extended to viscoelastic problems through the elastic-viscoelastic correspondence principle (9). Our moduli measurements were carried out at the -1

same strain rate (£ -0.02 min

) and therefore do not solve the general problan

of the moduli as a function of time. However, comparison between different experimental systems and theoretical models calculated taking into consideration the PE modulus measured at the same strain rate, can be useful. Experimental values of Young's tensile modulus E strongly depend on elongation. For example, Fig.1 shows the plot of the secant modulus E = a/e

vs. strain, for the system

PE-glass beads treated with SIC. All data relating to this work were obtained in the initial region at e = 0.1%. Fig.2 shows the E modulus of PE composites with glass beads untreated, or hydrolyzed and treated with different adhesion promoters: the modulus increases with the glass volume fraction C^ and depends appreciably on the type of adhesion promoter, increasing in the order:untreated glass (UT); stearylisocyanate treated glass (SIC); poly-co-acrylic

acid-

hexadecyl methacrylate treated glass (AA 90 - ECM 10); poly-co-vinylpyridinebutyl aerylate treated glass (VP 50 - AA 50). Analogous results were obtained when studying the flexural modulus of FE-beads treated with SIC or ITS (4). In parentheses are the monomer symbols and their content in weight percent.

Discussion

A large body of literature considers the composite as a two component system characterized by perfect bonding between spherical inclusions and natrix. The aim of the theories is to give an exact expression of the two effective elastic moduli (e.g. bulk and shear moduli, k and y) of the composite as a function of the corresponding moduli of the components and of their volume fraction. The models do not generally foresee a relation between moduli and surface treatment. Three component models, considering also the interface material, are nearer to the reality. We developed an elastic three component model

(10)

which is based on the stress and strain fields given by Matonis (11), for a tensile test, arid by Love (12) for an hydrostatic experiment, provided that an equivalent homogeneous medium surrounds the spherical inclusion, the interlayer

108

and the matrix. A system of twenty non linear equations with twenty unknowns can be written. No close-form solution has been developed as yet; only numerical solutions can be studied. For interlayer thickness going to zero the model reduces to the two component three phase model of Christensen (13). Matonis (11) analyzed the effect of an interphase on the stress distribution around an isolated inclusion and found that it has a amall influence on the stress field. We conclude that in the case of small or vanishing amounts of adhesion promoters two component rrodels can be used with good approximation. General elastic solutions which give the effective linear stiffness tensor C.., , can be found l^kl by solving Hooke's law to obtain < a . .> = C. 13 ljkl where

kl

< a .> and < e, ,> are the average stress and strain in the composite. kl

To perform the operation the exact solution of the stress and strain fields (a^j (X) and e ^ (X) is needed. Rigorous results are limited to a few idealized systems such as that of the dilute suspension spherical inclusion model (14,15). In this situation a single sphere is considered irrmersed in a matrix and the stress and strain fields can be rigorously determined, respectively, in a compression and in a shear ideal experiment; finally, by setting the strain energy in the suspension, equal to that of the equivalent homogeneous medium, the k and u moduli of the composite can be derived in an exact close form (Table 1) where indexes 1 and 2, respectively, refer to inclusion and matrix. The E modulus can be obtained using the equation E = 9 yk / (3 k + ij) . The Poisson ratio is v = (3 k - 2y ) /2(3k+u). The study of finite concentration has been seen to be much nore ccrtplicated and rigorous solutions have not yet been found. Consequently many attempts were made to delimit and construct upper and lower bounds for the moduli; energy methods were used for this purpose. The theorem of minimum potential elastic energy states that only the field that satisfies the equilibrium equations, among all the possible displacement fields, makes the potential energy an absolute minimum. Therefore, the actual strain energy in the equilibrated composite cannot exceed the energy of any fictitious non equilibrated state of

109

TABLE 1. ANALYTICAL EXPRESSION

TYPE MODEL SUSPENSION MDDEL

W

k = k2+

S

1+ ( k r k 2 ) / ( V

(14,15) JJ

ry2)

15(1-V2) (1-u/y C 2

U2

7-5V2+2(4-5V2) (Ui/y2) -1

REUSS and TOIGT bounds

k

(16)

,

k

~R

2

S C

1 k 1^2 k 2

-1

-L. H HASHES and SHTRIKMAN bounds (17)

y

A 2

s P *

k-k 1^ 2 (k r k 2 )/(k 2+ k e ) " k r k 2 " 1^ 2 (k r k 2 )/(k 2+ k u )

tok =

eT

V U

e

k

u= T"

y

1

= -L. (i. + 1 0 2 y2 9k2+8u2

3

/

1

if (Ul-y2) ( k ^ ) a 0

CHRISTENSEN Three phase model

k as in Hashin formala

(13)

A (y/y ) +2 B(y/u )4C = 0 For A,B,C see (13)

4.

10 ,-1

110 ANALYTICAL EXPRESSION

MODEL

(krk2)(4V3k2) S k = k + 2 4u 2+ 3k l+ 3(k 2 -k l )C 1

HASHIN bounds (18)

1+(1-u 1 /y 2 )y 1 c i for y^ and y^ see original paper

HALPIN TSAI model (20)

1+znC. 2

k

1_nC

1

v^ere Z = 4 ^ / ( 3 ^ ) ; W

V

2

U

C

W

V

1/U2+C_C1

( t J

y

n = 1

W

)

/u2"1)

where C = (7-5v2)/(8-10v2)

NIELSEN and LEWIS model (21)

E_ E„

1+ABC

1 1-tpBC.

where A = (7-5v2)/(8-1CK>2) B =

(E^-U/iE^+A)

1> = 1 + (1-CJC /cl m 1 m being C the maximum packing factor ( = 0.637 for m random close packing of spheres)

111

distortion having the same surface displacement. Assuming f o r example a l i n e a r l y varying displacement f i e l d , i . e . a uniform strain across a l l phases (one of the simplest but f i c t i t i o u s ) and decomposing the s t r a i n to a pure d i l a t a t i o n and a pure shear, one obtains, f o r the bulk and shear moduli u and k of the system, the conditions:

k £ k = E C.k. v 11

and

y < u = Z u.C. v 11

were v stays for Voigt, C

i s the volume concentration of the i canponent, k^

and vt are the bulk and shear moduli of the i component. Similarly i f we consider, among d i f f e r e n t loading arbitrary situations which a l l s a t i s f y the s t r e s s boundary conditions,a f i c t i t i o u s loading named of Reuss, giving uniform s t r e s s across a l l phases, i t w i l l not be in the minimum energy condition because i t does not s a t i s f y the compatibility equations, so that inclusions and matrix could not remain energy w i l l be higher than

1 A

£

1/k„

R

=

bonded. Therefore i t s complementary

that of the r e a l body; hence (16):

£ C . A .

1 1

and

1/y

i

1/u

R

=

£

C./\i.,

l

'l

where R stays f o r Reuss. The Voigt bounds always exceed the Reuss ones and correspond to a p a r a l l e l combination, while Reuss lower bounds correspond to a s e r i e s combination. These bounds, independent of the inclusion geometry, are unfortunately too f a r apart for t y p i c a l composites. Tighter bounds (Table 1 ) , also independent of the phase qeometry, can be obtained according to a procedure derived by Hashin and Shtrikman(17). These authors applied the c l a s s i c a l minimum principles of e l a s t i c energy, but considered a variable f i e l d of admissible s t r e s s and strain rather than the uniform f i e l d s used in the derivation of the Reuss and Voigt bounds. Another approach to the bounds of the moduli which r e f e r s to the particular spherical geometry of inclusions was made by Hashin (18) using energy variational theorems f o r an exact solution of the s t r e s s strain f i e l d s in a spherical composite. They should not be confused with the bounds of Hashin

112

and Shtrikman which are independent of the phase geometry. Hashin supposed that the composite is made by the size gradation of spherical particles characterized by r^/r2 = constant, where r^ is the inclusion radius and r^-r^ is the thickness of the embedding matrix. The upper and lower y bounds (Table 1) coincide at low and high volume concentrations C^of the inclusions. The y value at very small C^ is equal to the dilute suspension model expression. Upper and lower bulk moduli bounds coincide in all the concentration fields. A third class of interpretation of the gross moduli is based on schematic models and gives single value expressions. Solutions of the problem in the case of finite concentrations are numerous. We rartarber Christensen's composite spheres model (13) based on the existence of a spherical inclusion, an embedding matrix phase and an outer equivalent homogeneous medium. The resolving equations (Table 1) for the moduli coning frcm stress analysis and energetic considerations are exact for the model, but their applicability is conditioned by the reliability of the correspondence of the model to the real composite. Similar considerations can be made about the Kerner equations (19) which are based on a model similar to Christensen's composite sphere model, but with the interposition, between the matrix and the equivalent average material, of a layer which has properties shading continuously in unknown way from the matrix properties to the average medium ones: the stiffness tensor is obtained from the Hooke law after calculation of the average stress and strain. Finally, many empirical or semiempirical equations useful for practical application have been proposed. As an example, Table 1 gives the Halpinr-Tsai (20) equations which are an approximate form of Kerner's equation and the Nielsen-Lewis (21) equation which also derives from self-consistent models and takes into consideration the limited volume fraction that can be occupied by a particulate phase. Other empirical equations are related to equivalent mechanical models such as the well known Takayanagy models (22) which are convenient for gnpirical curve fitting but are not realistic. The Dobkcwski (23) parallel void fraction model in the filler space appears to be of seme interest for data fitting. Fig. 3 shows the E boundary values calculated according to

113

E (MPa) C = 0,1 V. 1200

_ O UT . x SIC _ • AA9C- EDM 10 . £ VP 50 - 8A 50

800

«00-

•a

200^ e(%) Fig.1

1

E l a s t i c modulus of composites PE-SIC treated g l a s s beads, at d i f f e r e n t f i l l e r voluire concentrations, versus strain.

0

c,

05

F i g . 2 Young ' s moduli versus f i l l e r concentration of PEg l a s s beads composite. For symbols see the t e x t .

HS»K

E=ar/. O UT X SIC a AA90-E0M10 & VP 50 -BA50

C = 0.17. o UT » SIC • A A90 - EDM 10 A VP50-BA 50 0 0.5 Ci Fig. 3 Voigt and Reuss (VR) model, Hashin - Shtrikman (HS) model, Kemer (K) model and experi= mental r e s u l t s (see f i g . 2 ) .

0.5 Fig.4

Moduli according: d i l u t e suspensions (DS), Hashin (H), Lewis - Nielsen (IN) , HalpinTsai (HT) models.

114

Voigt and Reuss (VR) and Hashin and Shtrikman (HS) using k = 467,000 2 k 2 = 18,800 y = 280,000 y = 760 In kg/cm . In the same figure we bring in comparison the E moduli of our composites. We can see that the lower H-S bound coincides with Kerner's (K) modulus. Fig. 4 shows the E modulus behaviour for other models: Hashin (H); dilute suspension (DS); Halpin-Tsai (H-T); LewisNielsen (L-N). We note that the experimental E data are always far from the parallel model predictions, and that no one of the presented moduli equations fits exactly the data. Therefore approximate fittings obtained by using the semi-empirical parameters appear to be still important, notwithstanding they cannot offer a true fundamental interpretation. Using the elasticity theory very few conclusive and exact results exist, based on a rigorous demonstration, able to predict the moduli of binary systems, such as the dilute suspension spherical particles theory, valid in a snail range of filler concentrations and the moduli bounds expressions which are, unfortunately,not very useful for practical cctrparison. Methods which consider particles and matrix embedded in a homogeneous material, equivalent to the composite (Christensen, Kerner), are interesting and theoretically rigorous, but do not adequately predict our experimental data, not only for the composites containing interface adhesives, but even for simple tvro component systems, like untreated beads-PE. The incomplete correspondence has to be attributed to a not completely true model. For example, the supposed homogeneity of the material beyond the matrix external limit, on which the self-consistent models are based, is a schematic unproven hypothesis. The presence of a third component (the adhesive) should also be considered for thin

interlayer, even if we do not expect large

changes in the moduli calculated according to the two-component models. Experimental data (Fig. 2) show, on the contrary, that modulus E can increase very appreciably according to the surface treatment. Another important hypothesis to be discussed, comon to almost all models used, is the perfect adhesion at the matrix-inclusion interface. Good bonding should be in principle easily reached in any practical case of molded composites due to the differential thermal contraction of glass and polymer. Nevertherless also in such practical cases the forrration of discontinuities, defects and small cavities at the interface is possible.

115 Table 2. Wetting Properties of the Interface A Glass-VP75 BA25 Polymer measured on plane sheet System adherent

PE VP-BA VP-BA

solid surface

HG OTG HG

Temperature of measurements

Contact angle

°C

150 205 205

50.4° 39.2° 27.8°

Polymer surface tension dyn/cm 17.8 17.4 17.4

Work of adhesion W= (1+cgse) dyn cm/cm

29.1 30.9 32.8

Symbols HG = hydrolyzed A glass; UTG = untreated A glass; PE = polyethylene LG1-1300 = nelt flow index 70 g/10'; VP 75 BA 25 = g copolymers of 4 vinylpyridine (75% by weight) and butylacrylate (25% by weight).

Fig. 5

Stress-strain curves at lew strain for different ccrrposites (see text).

116 Defects induce stress field variation and act as stress concentrators with consequent decrease in the elastic modulus in a way that current theories cannot predict. Defects could be related to the surface wetting conditions which depend on the wetting driving force and kinetics. Table 2 shows, as an example, the contact angle, the surface tension and the vrork of adhesion between PE and glass (as a reference) and between the grafted VP-BA copolymer and glass in order of increasing adhesion work

W. Fig. 5 shews the

behaviour at low elongations (elongation rate = 1 nm/min) of composites EE 20% volume glass beads untreated or treated with

4vinylpyridine butylacrylate

copolymer (1% on the glass weight). The modulus appears to increase with increasing W. The importance of preliminary glass hydrolysis can be seen from the data. The good affinity of VP-BA put on the interface for PE can be derived from the very low contact angle of hydrocarbons on polybutyl acrylate homopolymar ( < 5° with n-octane). Good wetting favours the particle dispersion reducing beads agglomeration responsible for inhcrnogeneities which can also influence the modulus (24). Failure The presence of rigid spherical inclusions normally degrades the strength of the matrix material, because of the stress concentrations around the particles. Fig. 2 of one of our preceding papers (4) clearly shows this fact in the case of PE-glass untreated beads composite. Yielding happens at each glass concentration. Following Smith (25), Nicolais and Narkis (26), it is possible to foresee a lower theoretical limit of to the polymer only. Vfe get

/

a°

ay

attributing the loading capacity 2/3 = 1-iC^/C^) where a° is the

polymer yield stress, C the beads concentration at the maximum packing factor, m and °Y

Oy 0

^ Y

the theoretical limit. We verified in all our PE composites that l ar 9 e ly depends on the surface treatment and is always well

above

the theoretical ratio. For hydrolyzed glass concentration C^ =0.2, using the VP 50-BA 50 and VP 75.5-BA 24.5 copolymers as adhesion promoters we obtained Oy

/ cr

=1.89 and 1.36, respectively. With the aim of finding more

fundamental insight, we shall now concentrate our attention just on the initial failure of the composite: for this purpose we prepared a glass beads-polymer

117 composite, molding two EE films and introducing between them a few beads, untreated or treated with 1% of VP 75.5 RA. 24.5 copolymer, and remolding. The system is theoretically comparable to the dilute suspension model. Observing the samples through an optical microscope during the tensile test, we found that failure always begins in the polar zone both with treated

and untreated

beads, and extends giving origin to dewetting and formation of vacuum ellipsoids differently developed in the tvro hemispheres. Fig. 6 shows successive moments of the failure. Looking with SIM at the fracture final section of the composites we saw that beads, when treated with VP-BA adhesives, are completely covered by the polymer. These observations extend the validity of our previous work (4). We conclude that in the

presence of adhesives,

failure happens in the matrix, beginning at the pole with reference to the tensile force direction. This allows us to apply the Goodier stress analysis (27) for the matrix only. Failure mechanism can be generally reduced to micro or macro cavities formation in the literature to foresee

and/or shear yielding. Many criteria are proposed the failure when reaching a critical value

alternatively of the highest principal stress (0 ) ; principal strain (ej); maximum shear stress ( ,

dilatation (A); strain energy (VM; shear strain

energy (WQ). Using the formulas of Goodier (27) and the expression of the principal stresses and strains (12) vie have calculated all the above defined functions for the conditions r/r = 1 where r is the inclusion radius and r o o is the distance from the center, and for G £ 90 that is to say, at the interface and for différent angles. Fig. 7 shows for example the ratios: o

/S, e /S, A E /S, T /S where S is the tensile external stress; analogously, m 1 2 1 2 W E /S and W E /S (where E is the matrix modulus) can also be represented, s m D m m Considering that failure begins in the polar zone, we can exclude as possible causes of failure all

the stress and energy functions which present maxima

far frcm 6 = 0. There remains to be considered 6

the maximum dilatation at

= 0, or the maximum principal stress, o , which shows a slowly varying

behaviour at

6 below 20°. The Goodier analysis does not foresee maxima of

the above considered functions for r > ro and 0 around zero, higher than the values found at the interface with the exception of

a ;

a^/S has in fact an

absolute maximum at 9 = 0 and r/r = 1.15. These considerations confirm that

118

aj or

A are responsible for the failure. The stresses of thermal origin,

t

orr' , a 66 ™ , o *#

due to the differential contraction of glass arid PE frcm

the molding to the room temperature, must be added to the Goodier a

0

a

rr»

17633 0113

6 0' i|)ijj exP

!

- The radial contribution, o ^ , is a ccrrpression stress

(negative) while 96

= a\p\j) =

-arr 'r~J2 are tensile contributions. The

thermal shear stress is zero. It can easily be demonstrated that dilatation induced by the external forces is not changed by thermal stresses, while the maximum principal stress

a^ is influenced by

a^T , but not considerably

at the level of the external force S that we found necessary for the failure. We conclude that failure in the case of adhesive beads is ruled by A nax or Oj rrax. A different situation appears to be met in absence of chemical or chanicalphysical adhesion, when thermal differential contraction between glass bead and polymer is the cnly responsible for adhesion. We observed by SEM in fracture sections that untreated beads are not significantly covered by adherent polymer. In this case we believe that failure begins just at the interface in the polar zone where arr is maximum, when arr caused by external forces equals A due to thermal compression. On the surface OQQ and o are not active and 6 = 0 for 6 = 0 . Therefore, in absence of W r6 adhesion promoters, detachment occurs at S values lower than in their presence. Vfe could measure the increase in volume due to vacuum formation of different samples during the tensile tests, by contanporary measurements of elongation and lateral size variation of the samples and, also, by independent measurement of density. We observed that the increase in volume begins at a specific elongation of about 1.8% in the case of composites with untreated beads, and at a higher strain ( e

=3%) with hydrolyzed beads treated with VP 74.5 BA

45.5 while PE alone does not show volume increment. Qualitatively this is in complete agreement with the theoretical predictions.

119

Fig.6

Sequence of cavities formation, during the tensile test

o,/S

Fig.7

Highest principal stress a normalized to external stress S; normalized highest principal strain f^E^/S, normalized maximum shear stress T /S,normalized maximum dilatation AE /S. I m

120

Acknowledgement The financial aid of CNR (Progetto finalizzato chimica fine e secondaria) and of Ministero Pubblica Istruzione is gratefully acknowledged. Thanks also to Prof. A.Penati, Prof. F.Severini, Dr. L.Di Landro and many students for their contribution to this vrork.

References 1. Christensen R.M.. 1979. Mechanics of composite materials. J.Wiley N.York, p.7. 2. Pegoraro M., Pagani G., Clerici P., Alessandrini G., Maggione R.. 1974. Ing. Chim. Ital. 161. 3. Pegoraro M., Pagani G., Clerici P., Penati A.. 1977. Fiber Sci. and Techn. 1£, 263. 4. Pegoraro M., Penati A., Canmarata E., Aliverti M. 1984. In Polymer Blends Vol.2. (Kryszewski M., Galeski A., Martuscelli E., eds). Plenum Press N.York, p. 205. 5. Mistrovica V.. 1983. Copolymers of acrylic acid and hexadecylmethacrylate. Thesis at the Polymer Science School G.Natta. Polytechnic of Milan. 6. Penati A., Meregalli M., Pegoraro M.. 1984. Riv.Staz.Sperim.Vetro, 11. 7. Severini F., Pegoraro M., Saia L.. 1985.Atti Conv.Ital.Chem.Soc. Siena C.I.10 8. Neumann A.W., Good R.J. 1979. In Surface and Colloid Science, Vol. 11 (Good R J, Stromberg R.R., Eds) Plenum Press N.York, p.31. 9. Schapery R.A. 1974. In Mechanics of Composite materials (Sendeckyied G.P.) Academic Press N.York, p.85. 10. Garavaglia F.. Moduli elastici dei conpositi a microsfere. Thesis at the Polytechnic of Milan. 11. Matonis V.A., Small N.C.. 1969. Polymer and Engin.Sci. 9, 100. 12. Love A.E.H.. 1944. A Treatise on the mathematical theory of elasticity. Dover Publications.

121

13. Christensen R.M.: reference 1,p.52. 14. Christensen R.M.: reference 1,p.41. 15. Rizzi A., Coppi A.. 1984. Moduli elastici di materiali compositi particellati. Thesis at the Polytechnic of Milan. 16. Paul B.. 1960. Trans. A.I.M.E. 218, 36. 17. Hashin Z., Shtrikman S.. 1963. J.Msch.Phys.Solids 11, 127. 18. Hashin Z.. 1962. J.Appl.Mech. 29, 143. 19. Kerner E.H.. 1956. Proc.Phys.Soc. 69 B, 808. 20. Halpin J.C., Kardos J.L.. 1976. Polymer Eng. and Sci. 16/ 344. 21. Lewis T.B., Nielsen L.E.. 1970. J.Appl.Polym.Sci. 14/ 1449. 22. Takayanagy M., Harima H., Iwata Y.. 1963. J.Soc.Mater.Sci. Jpn 12, 389. 23. Pegoraro M., Dobkowski Z., Di Landro L., Garavaglia F., Penati A.. 1985. Annali di Chimica 75, 223. 24. Nielsen L.E.. 1974. Mechanical Properties of Polymers and Composites. M.Dekker N.York. 25. Smith T.L.. 1959. Trans.Soc. of Rheology J3,113. 26. Nioolais L., Narkis M.. 1971. Polym.Eng. and Sci. 11,195. 27. Goodier N.J.. 1939. J.Appl.Mech. 55, 39.

OXIDATIVE DEGRADATION OF POLYPROPYLENE CATALYSED BY MINERAL FILLERS

K.Vesely,

O.Petrfij,

Chemopetrol 656 49 Brno

A.ZahradniCkovi

Research

Institute

of Macromolecular

Chemistry,

Int roduct ion The m u l t i d i s c i p l l n a r y involves latter

t o p i c concerning the composite

but a l s o

the f i l l e r s compounding, application

the k i n e t i c s

act

The o x i d a t i v e

as

it

impossible

effect

it

reason,

p o i s o n i n g metals. content

of

contact

certain metallic

of

polymer material

compounds namely

rubber producers

the chemical p u r i t y of

however,

shown that

to the c a t a l y t i c of

individual the p u r i t y

effect.

Fe^O^. Not only that

effects light

transition

This this

during the p o l y o l e f i n

defined is

metals.

rubber-

fillers

e.g.

by the

The p r a c t i c e

in t h i s way i s not

clearly

illustrated

compound e x h i b i t s

those

for a long

these metals are sometimes c a l l e d

So f a r ,

of

d u r i n g the to preclude

application.

micro-ground calcium carbonate has been c h a r a c t e r i z e d overall

which

during

might be p o s s i b l e

to prevent

of Cu, Mn, Fe has been known to the For t h i s

the

during p r o c e s s i n g by working in an i n e r t

with oxygen during p r a c t i c a l The d e t r i m e n t a l

at

p r o c e s s i n g and predominantly

Although

plastics

is

reactions

i n p r a c t i c e proceeds a l r e a d y

thermoplastic of a r t i c l e s .

atmosphere,

of d e a r a d a t i v e

of

catalysts.

degradation

the o x i d a t i o n of

time.

materials

not only the s t r u c t u r a l and mechanical s t u d i e s

zero

degradation but a l s o

has,

related

on example catalytic

it

a c t s as a

stabilizer.

The compounds c o n t a i n i n g a crucial

transition

r o l e during a u t o o x l d a t i o n

metals have been known to

play

reactions.

of

The importance

these compounds as oxygen a c t i v a t o r s was noted as e a r l y by George and Robertson (1) and l a t e r The process was d e s c r i b e d by the Mn+.

02

+

RH

>

as In 1946

on a l s o by other authors

reaction:

M n+ + R* + H0 2 '

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

(2,3).

124

On the o t h e r hand Lun£k a t a l . mechanism of c a t a l y t i c

(4,5)

a u t o o x i d a t i o n based on oxygen

i n t o t h e C-H bond. T h i s process t r a n s i t i o n m e t a l complex state.

have r e c e n t l y put

insertion

i s b e l i e v e d f o be c a t a l y z e d by a

i n which t h e complex i s

The f o l l o w i n g c a t a l y t i c

f o r w a r d the

i n a l o w e r valency

c y c l e has been suggested by these

autors: Mn+ • 0 2

Mn+.02

»

M n + . 0 2 + RH

c e n t r e M n + i s formed through t h e r e d u c t i o n of

where t h e c a t a l y t i c M

n+1

.

ally.

ROOH + M n +

>

T h i s r e d u c t i o n may be induced e i t h e r

t h e r m a l l y or photochemio-

One of t h e main e v i d e n c e s u p p o r t i n g t h i s mechanism i s seen

i n the absence of R 0 2 .

recombination

products.

According t o our o p i n i o n , when c o n s i d e r i n g the a u t o o x i d a t i o n

pro-

cess c a t a l y s e d by t r a n s i t i o n m e t a l compounds i t

to

assume the involvement ligand f i e l d .

of

radicals

is p l a u s i b l e

t h a t a r e complex bonded i n

formed i n the r e a c t i o n of t - b u t y l h y d r o p e r o x i d e w i t h and a c e t y l a c e t o n a t e s ( 7 ) on a s e r i e s of bonded r a d i c a l s .

naphtenates

of manganese, c o b a l t or vanadium was described

as e a r l y as i n 1967 by Brandon and E l l i o t

(6).

I n 1971 we reported

r e a c t i o n s e x h i b i t e d by these l i g a n d

field

The t o p i c c o n c e r n i n g complex bonded r a d i c a l s

been r e c e n t l y a s u b j e c t

of c o n s i d e r a b l e

E . S . G o u l d has been p u b l i s h e d One of the v e r y i m p o r t a n t

interest

(8). autooxidation

to

inhibitor

which i s observed when c e r t a i n c o n c e n t r a t i o n of t r a n s i t i o n Black

metals f o form complexes w i t h h y d r o p e r o x i d e s . c o n s i d e r e d by the a u t o r as a c a t a l y s t ium c a r b o n a t e f i l l e r s out e x p e r i m e n t s e i t h e r matrix. at

the c a t a l y t i c

of

T h i s complex

of

is

oxidation.

i n f l u e n c e of v a r i o u s

on p o l y p r o p y l e n e o x i d a t i o n ,

we have

calc-

carried

i n an l i q u i d model or in a p o l y p r o p y l e n e

K i n e t i c measurements i n l i q u i d phase a r e p o s s i b l e

relatively

metal

( 9 ) has accounted f o r t h i s by t h e a b i l i t y

In order to c h a r a c t e r i z e

has

and a r e v i e w by

phenomena in the f i e l d of

i n non p o l a r media i s t h e c o n v e r s i o n of c a t a l y s t i s exceeded.

the

The f o r m a t i o n of l o n g l i v i n g complex bonded radicals

low t e m p e r a t u r e s i n c e the r a t e c o n t r o l l i n g

i n t h i s system i s oxygen d i f f u s i o n even a t

intense

only step

stirring.

125

Influence of Calcium Carbonate Impurities on Trlmethylcyclohexane Oxidat ion Trlmethylcyclohexane (TMC) was used as a liquid model of polypropylene chain. The oxidation was carried out in a stirred reactor at 60°C, the oxidation rate being followed by the oxygen uptake method (10). Two samples of calcium carbonate were examined in basic runs. The principle characteristics of these fillers are presented in Table 1. Table 1.

Characteristics of Micro-ground Calcium Carbonate Samples

Denotat ion

Spec.surface

Content of Fe 2°3 ppm

?

m /g A

3.2

B

1 .4

Content of MnO ppm 16 41

53 143

15 ¿0 2 [mol.l" 1 ] x103

10

5

50

r

100

T i m e [min]

Fig.l.

150

Time dependence of oxygen uptake during TMC oxidation. 1 - Calcium carbonate A: lg/5 ml TMC. 2,3,4 - Calcium carbonate B: 0.4; 1.0 and 1.4g/5ml TMC.

126 The c o u r s e of o x y g e n a b s o r p t i o n

in the p r e s e n c e of b o t h

calcium

c a r b o n a t e t y p e s is s h o w n in F i g . l . It follows that the less c o n t a m i n a t e d c a l c i u m c a r b o n a t e A exhibits an i n d u c t i o n p e r i o d , the s u b s e q u e n t o x i d a t i o n b e i n g s l o w e r t h o u g h the s a m p l e has a g r e a t e r s p e c i f i c s u r f a c e .

If w e

even

consider

the s l o p e of the c u r v e a f t e r the e n d of i n d u c t i o n p e r i o d as a c r i t e r i o n of the r e a c t i o n rate it is s e e n that the T M C rate is p r o p o r t i o n a l to the s u r f a c e of b o t h c a l c i u m

oxidation

carbonate

types(Fig. 2).

2

U S [m 2 1" 1 ] *10"2

Fig.2.

D e p e n d e n c e of i n i t i a l rate (v) on the s u r f a c e of c a r b o n a t e s A a n d B. A C h ^ O H - 0.

calcium

M o r e o v e r , w e h a v e f o u n d that the p r e s e n c e of w a t e r has a s i g n i f icant c o - c a t a l y t i c

I n f l u e n c e w h i c h is c h a r a c t e r i z e d

In

Fig.3.

127

-4

Fig.3.

-3

-2

-1

J)

1 [H20]

3

4

Dependence of oxygen uptake rate (v) on the amount of added w a t e r ACH^OD . Calcium carbonate B lg/5ml TMC.

It is known that simple drying of the filler at 100°C is not sufficient to remove the sorbed w a t e r . It seems that in an ideally non-aqueous medium the oxidation does not take place. Hudec

(11)

found that cobalt complexes, known as oxygen absorption agents were functioning only when trace concentrations of w a t e r were present. This effect of water is seen in the influence of the latter on the symmetry of the complex

formed.

The influence of iron and manganese compounds on TMC oxidation in the presence of calcium carbonate was examined by

following

the oxidation of substrate in which calcium carbonate A impregnated by

the above mentioned metallic compounds was used as catalyst.

As shown in Fig.4., manganese exhibits a much higher effect than

catalytic

iron.

Moreover we have found that the catalytic activity is strongly influenced by the character of an anion bonded to the

transition

metal ion. The following sequence of anion catalytic activity was found for a series of manganese CI"

»

SO

4

2-

>

NO3"

^

compounds: CO

2-

3

128

Fig.4.

Dependence of oxygen uptake initial rate on MnCl- concentration (impregnated calcium carbonate A) o. Dependence of stationary rate of oxygen uptake on FeCl., content (impregnated calcium carbonate A) J. lg of treated CaCO_ i per 5ml TMC.

Polypropylene Oxidation Filled polypropylene samples were prepared in form of 0.06-0.1 mm thick seets by mold-pressing of a mixture compounded on laboratory kneader at 220°C. The samples contained 40% CaCOg, 0.3% stearic acid and 0.15% BHT. The reference measurement was carried out with non-filled polypropylene. For this purpose a commercial biaxially oriented sheet containing 0.125% BHT, 0,15% AO 49 and 0.15% DSTDP was used. The antioxidants were removed from polymer by extraction with hexane-chloroform-ethanol mixture. The hexane component swells the polymer while other solvents dissolve the additives. It was found that after 24 hours extraction the original amount of antioxidants was less than 4 p.p.m. The samples were placed into an oven with forced circulation of air. Within the sample thickness from 50-100 ¿im and up to 125°C, the rats of oxidation was independent of sample thickness i.e. the diffusion

129 was not a l i m i t i n g

process.

The oxidation k i n e t i c s were followed by monitoring the hydroperoxide group concentration.

The method due to PetrOj and Marchal

employed for the determination which i s based on f e r r i c spectrocolorimetry.

In order to achieve q u a n t i t a t i v e

the sheets were f i r s t adding

left

the reagents.

to swell

the f e r r i c

thiocyanate

512.5 nm. The method lenables us to determine

hydroperoxide concentrations with an The described method determines peroxyaclds.

reaction,

in benzene f o r 24 hours before

The absorbance of

complex was measured at

(12) was

thiocyanate

accuracy of -

lxl0~ 6 mol /kg.

the sum of hydroperoxides and

the l a t t e r being formed by o x i d a t i o n of side CH3~

groups. By adding diphenylsulphide

the peroxyacids are

decomposed within several minutes while the concentration remains unchanged.

selectively

hydroperoxides

In t h i s way, the peroxyacid

concentration may be determined from the d i f f e r e n c e of two analyses. In the f o l l o w i n g text

the r e s u l t s are presented of

measurements of polypropylene -filled)

kinetic

(both f i l l e d with 40% CaCOg and non-

thermooxidetion c a r r i e d out at 110 and 120°C, resp. The

calcium carbonate f i l l e r s Table 2.

are s p e c i f i e d

in Table 2.

CaC0„ Sample C h a r a c t e r i s t i c s

Denotat ion

Producer

Fe 2 0 3

%

MnO

Si0 2 +Al 2 0 3

%

Spec, surface

%

m 2 /g

64-Durcal 2

Pluees-Staufer

0.020

0.0015

0. 20

140

PobSiovice

0.057

0. 006

2. 28

2.4

0.008

0.58

4.6

0.006

0.40

4.3

II

155

Pob82ovice I

0,061

127

ÜNS Kutn4 Hore

0.035

The k i n e t i c

curves

r e f l e c t i n g the hydroperoxide

dependence e x h i b i t , character. ates

after

induction p e r i o d ,

3.3

concentration-time

an a c c e l e r a t i n g

An example of measurements at 120°C in l o g - l o g coordin-

i s presented in

Fig.5.

130

log [ROOH] -2

•3

0

Fig.5.

0.5

log t (h )

1.5

Course of kinetic curves in log-log coordinates

It follows from Fig.5

(120°C^.

that the reaction order is differing

two periods, this being indicative of two different

for

reaction

mechanisms operating in each period. When the kinetic curves are plotted in the semilogarithmic

scale

(i.e. time dependence of log CROOHJ) than the straight lines are obtained which, in the case of non-filled sample and purer fillers (No 64 and 127), exhibit different slopes in both regions w h i l e in the case of less pure fillers

(No 140 and 155) only one slope

is observed as it follows from Fig.6

(at 120°C).

W h e n the catalytic activity of a given filler is characterized according to the slope of log ROOH - t i m e dependence then for given fillers the values shown in Table 3

are obtained.

131

Table 3.

Catalytic Activity of Fillers (40% in PP)

Filler

Slope

Non-filled polypropylene

( L O ^ S " 1 ) at temperature 110°C 120°C -

1.3 4.5

64 - Durcal

0.7

127 - GNS Kutn6 Hora

1.7

7.7

155 - PobSiovice I

5.3

13.2

140 - PobSiovice II

6.6

14.8

Fig.6.

Course of kinetic curves in logCROOHH-t coordinates (120^.

The catalytic activity of fillers is directly proportional to manganese concentration on the filler surface as it follows from Fig.7.

132

5

1

Fig.7.

2 M n O / S [g.m 2 ]*10 5

Dependence of filled (40% CaCOg) polypropylene oxidation rate upon the surface concentration of manganese.

Discuss ion The employed experimental technique enables one to characterize the overall catalytic activity of fillers. It follows from Table 3 and Fig.7 that the latter property increases with the manganese content. There has not been enough of results allowing a more detailed elucidation of elementary reactions taking place during the filler catalysed autooxidation. For this reason we restrict ourselves to discussion of plausible alternatives. It is assumed that the first - slower - phase of oxidation is a catalytic process taking place within the ligand field of metal catalysts; in the second phase we assume an outer sphere radical react ion. First catalytic phase. Lunik et al. (5) have suggested that the initial hydroperoxides are formed through the insertion reaction of activated oxygen - Cat.0 2 - into the R-H bond, the process taking place in the ligand field. In such a case the hydroperoxide

133 increase should be independent of its concentration

i.e. the

following relationship should hold: CROOH: » k.t It seeiss that this relation is valid only for the very reaction stages, provided that metal concentration

initial

is low. As soon

as a certain minimal hydroperoxide concentration is reached a complex Cat.ROOH is likely to be formed which is, according to Black

(9) relatively stable, i.e. the concentration of free ROOH

will be very low. The rate determining step in initiation

process

w o u l d thus be the transformation of the mentioned complex

to

complex bonded C a t . R 0 2 *

radicals. If one assumes that the propag-

ation reaction in the ligand field proceeds relatively

easily

(see ref.7) then the high value of kinetic chain length may be accounted for. Taking into consideration the findings of authors (ref.4,5) who have not observed any products of radical

recombin-

ation it seems likely that the disappearance of complex

bonded

radicals takes place directly in the ligand field (in the vicinity of filler) through the reaction with filler or w i t h the present metal which is not identical with the active centre. As shown by Walling

(13), the metallic compounds in the higher valency state

may act as efficient radical scavengers. This effectively that the number of centers on w h i c h RO^' radicals may will be considerably higher than the number of

means

disappear

catBlytically

active centres CCat] . When formulating the termination reaction in the ligand field as k

R0 2 * + S

t1

—>

products

(S being the termination active filler surface) then for the stationary state the following expression for CR0 2 ']

will be

valid: dCROg*] p

CR00H]

_ k t CR0 2 'D CS3

= 0

dt W h e n this expression is Inserted into the relation of propagation dIROOH]_

k

P

CRO

2

CRHD

134 which upon integration yields: InCROOH:«

k t cs:

• t

This relationship corresponds to the experimentally observed results. An alternative Interpretation of the found semllogarithmlc dependence of hydroperoxide concentration on time is based on the notion that both the first and the second oxidation phase proceeds, in the case of catalytically active fillers, from the very beginning i.e. there is an absence of induction period. Second - fast - oxidation stage. Upon achieving a certain critical concentration of hydroperoxides a faster initiation reaction will become more pronounced which can be written as follows: 2 ROOH.Cat

>

RO* + Cat.R0 2 * + H 2 0

Fres RO* radicals initiate the chain reaction in polymer matrix. The disappearance of radicals through the reaction with filler"' surface will be considerably more difficult. For this reason the blnolecular radical recombination is believed to be more probable. The over-all process may be described by the following equation: d CROOH3 dt

1/2 -

k

i

neat:

crooh:

» k.CROOHJ

t /

which gives upon integration crooh:

» k

0

It follows that this squat ion holds for both filled and non-filled samples as documented in Fig.8 where the time scale was shifted to coincide with the beglning of the second autooxldation phase.

135

Pig.8.

Time dependence of s e c o n d p h a s e o f a u t o o x i d a t i o n log-log coordinates.

The h i g h o v e r a l l

reaction

r a t e at

rate

in

t h i s s t a g e may be e x p l a i n e d

by

a r a t h e r s l o w r e c o m b i n a t i o n p r o c e s s o f RO^" matrix

(Buchachenko)

recombination magnitude

r a t e of

in atactic

( 1 4 ) . So e . g .

Niki

radicals

in

polymer

(15) has reported that

RO,,* r a d i c a l s was s l o w e r by 3 - 4 o r d e r s p o l y p r o p y l e n e when compared t o the

the

of

liquid

phase. It

i s worth mentioning that p o l y p r o p y l e n e

relatively after

stable

even a t

_2

r e a c h i n g c o n c e n t r a t i o n of

able decomposition occurs so that increasing. carbonyl

ca 4 x 1 0 their

mol/1 t h e i r concentration

are Only

consideri s no longer

The main p r o d u c t o f h y d r o p e r o x i d e d e c o m p o s i t i o n

are

c o n t a i n i n g compounds whose f o r m a t i o n c o u l d be m o n i t o r e d

by I R s p e c t r o s c o p y

i n the c a s e o f n o n - f i l l e d

The f o l l o w i n g mechanism h a s been put (16) to account

polypropylene.

f o r w a r d by P e t r O j and Mardial

f o r the f o r m a t i o n o f c a r b o n y l s k

R O2 *

hydroperoxides

temperatures as h i g h as 120°C.

+ wCH-v» I 00H

c

^

ROOH + - w - C - w II 0

in

polyethylene:

+ HO*

136 A selective reaction of R0 2 *

radicals with tertiary hydrogen on

ROOH is believed to be preferred as a consequence of its activation due to the presence of vicinal polar groups.

Conclus ion A detailed study of filler chemical reactivity is considered necessary in connect ion with further development of

composite

materials. Both the interfacial structures and the character of matrix degradation in the course of processing and practical application should be examined. The study of first phases of catalytic oxidation in non-stabilised systems is believed to be of primary

importance.

References 1. George, P., A.Robertson. 1946. Trans.Faraday Soc. 42, 217. 2. Wittig.A. 1948. A n g e w . C h e m . 60, 169. 3. Uri.N. 1956. Nature 177, 1177. 4. Lunik,S., P.Lederer, F.Stopka, 0.Vepfek-SiSka. Czech.Chem.Commun. 46, 2455.

1981. Coll.

5. Lunik,S,, 0.VepPek-Si§ka. 1984. 16th Symposium on Catalysis, Liblice. 6. Brandon,R.W., C.S.Elliot. 1967. Tetrahedron Lett. 4375. 7. TkiC.A., K.Vesely, L.Omelka. 1971. O.Phys.Chem. 75, 2575, 2580. 8. Gould,E.S. 1985. Acc.Chem.Res. 9. Black,3.F. 1978. O.Amer.Chem.Soc.

22. 100. 527.

10. ZahradniCkovi.A., K.Vesely. 1982. Chemicky prOmysl 32, 533. 11. Hudec.P. 1978. 3.of Catalysis 53, 228. 12. PetrOj.O., O.Marchal, S.Zehnacker. 1980. 14th French-Czechoslovak Conference on Oxidative A g i n g and Burning of Polymers. 13. Walling,C. 1975.

Acc.Chem.Res. 8, 125.

14. Buchachenko.A.L.

1976. O.Polym.Sci., Symposium No.57, 299.

137 1 5 . t N i k i , E . , C.Decker, 11, 2813. 16.

PetrOj.O.,

F.R.Mayo.

O.Marchal.

1980.

1973. 0 . P o l y m . S c i . , Radiat.Phya.Chem.

16,

Pol.Chem.Ed. 27.

PREPARATION-PROPERTIES-USE 1. Polyolefin Composites

141-250

2. Other Polymer Composites

251-328

3. Thermosetting M a t r i c e s

329-398

SOME PROPERTIES OF HIGHLY FILLED ORIENTED POLYOLEFINES WITH CHALK AND OTHER SYSTEMS

M.Kryszewski, A.Gal^ski and T.Kowalewski Centre of Molecular and Macromolecular Studies Polish Academy of Sciences, 90-362 Lddf, Poland

Introduction The use of fillers in semicrystalline thermoplastics appears to have a huge importance. The application of such materials ranges from simple filling to more elaborate structures like laminates, chemically modified and oriented systems. The very different fillers which can be used, e.g. minerals and cellulosic materials, make possible to obtain composites with various properties. The.» large number of literature data and results of systematic studies in various areas of filled thermoplastics will be not quoted here. The mechanical parameters obtained from stress-strain measurements (modulus, breaking stress and strain) and impact strength, modulus of shrinkage and water sorption investigations make possible to evaluate the role of fillers in different composites. Chalk an^ other anisometric fillers are of particular interest because they make possible to obtain composites with different properties in two dimensions after large deformation. Actually there is a need for low density composites which are characteristic as well of high impact strenght and flexibility as of sorption of water and other liquids. These features may be obtained by controlled deformation of composite samples leading to orientation of polymeric matrix and formation of voids around the filler particles. Such materials should exhibit good sorption properties and flexibility perpendicular to the deformation direction. It will be shown that the proper use of fillers can be of primary importance to achieve such goals. The high degree of orientation of the polymer matrix and of filler particles by cold drawing is possible only when the polymer-filler adhesion is not too strong which prevents the premature formation of cracks

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

142 (note t h a t a l l w i d e l y a p p l i e d f i l l e r t r e a t m e n t s , e . g . use of s i l a n e s , are designed t o i n c r e a s e the a d h e s i o n ) . The nature of chalk and p a r t i c u l a r l y of wood based f i l l e r s does not promote the adhesion t o o l e f i n i c t h e r m o p l a s t i c s .

Various

i n v e s t i g a t i o n s have shown, however, t h a t the a b i l i t y f o r deformation of p o l y m e r - f i l l e r

plastic

i n t e r f a c e in such composites

is

r a t h e r low (low impact s t r e n g t h and e l o n g a t i o n a t break in comparison with pure p o l y m e r ) . Mineral f i l l e r s c r y s t a l l i z a t i o n of the polymer matrix

[lj,

i n f l u e n c e on the may introduce

geometri-

c a l d e f e c t s a c t i n g as notches, and due t o agglomeration may cause t o o strong s t r e s s i n h o m o g e n e i t i e s . T h e r e f o r e , according

to

our p r e v i o u s s t u d i e s chalk was t r e a t e d w i t h i n t e r f a c e m o d i f y i n g a g e n t . The l i q u i d l a y e r

( o l i g o m e r of e t h y l e n e o x i d e M w =300) around

chalk p a r t i c l e s has been shown t o i n h i b i t the f r a c t u r e phenomena. This l e d composites

t o high impact s t r e n g t h and e l o n g a t i o n a t break of

the

^2-4] and made p o s s i b l e c o l d drawing and o r i e n t a t i o n

of h i g h l y f i l l e d

samples t o high draw r a t i o

(X=5 t 6)

[5,6],

The aim of t h i s paper i s t o d i s c u s s some r e s u l t s of our on the e f f e c t of high f i l l i n g

studies

and o r i e n t a t i o n on morphology and

s o r p t i o n p r o p e r t i e s of composites obtained from m o d i f i e d chalk and i s o t a c t i c p o l y p r o p y l e n e

( i - P P ) . These r e s u l t s a r e compared

w i t h those o b t a i n e d when studying composites c o n s i s t i n g of the same polymers w i t h wood based f i l l e r s without s p e c i f i c s u r f a c e

treat-

ment. The p o s s i b l e r o l e of i n t e r f a c i a l

cellulo-

agent in the case of

se f i l l e r s w i l l be i n d i c a t e d c o n s i d e r i n g

s e v e r a l c o n d i t i o n s which

have t o be f u l f i l l e d in order t o reach m a t e r i a l s of comparable w i t h those of o r i e n t e d

polyolefins

properties

highly f i l l e d

with

modified chalk.

Mineral F i l l e r s The e f f e c t of

in S e m i c r y s t a l l i n e

liquid

Thermoplastics

i n t e r f a c i a l agent on mechanical

properties

of the composites has been a l r e a d y shown f o r p o l y p r o p y l e n e and p o l y e t h y l e n e s f i l l e d w i t h c h a l k , t a l c and k a o l i n . The composite samples e x h i b i t t y p i c a l

s e m i c r y s t a l l i n e thermoplasic

behaviour:

necking and p l a s t i c d e f o r m a t i o n . Chalk f i l l e d p o l y p r o p y l e n e has been taken as an example d e s p i t e some s p e c i f i c d i f f e r e n c e s found f o r composites w i t h t a l c and k a o l i n .

143

Extensive studies on the structure and deformation mechanism of polypropylene filled with chalk modified with oligomer of ethylene oxide have been carried our in our laboratory in recent years. Their results have been described in details in another papers. Below we briefly summarize some of them, important from the point of view of the present work. Observations of the uniaxial deformation of i-PP filled with various amounts of modified chalk have been carried out on INSTRON tensile testing machine equipped with the video system. Volume changes of marked sections of oar-shaped samples (0.4 mm thick, 5 mm wide, 10 mm long) deformed at a rate of 2 mm/min, were recorded as a function of a draw ratio measured between the adjacent marks. Results for the compositions containing 10,30,40, 50,60 wt.% of modified chalk (10 wt.% of OEO in relation to chalk) are shown in Fig.l. These results clearly indicate formation of

2

3 DRAW

Fig.l.

4

5

6

RATIO

Relative increase of volume vs. draw ratio during uniaxial deformation of filled i-PP with various amounts of modified chalk.

144 v o i d s in t h e s a m p l e s d u r i n g d e f o r m a t i o n . i n c r e a s e s w i t h t h e amount o f a f i l l e r filler

i s observed. This e f f e c t

d e f o r m a t i o n mechanism ( f o r d e t a i l s [6]),

ratio

i.e.

contents

( w i d t h 1 . 2 ym,

aspect

6) f i l l e r p a r t i c l e s a l o n g t h e d e f o r m a t i o n d i r e c t i o n

concentration

in l e n g t h .

On t h e m a c r o s c o p i c l e v e l

sample e x h i b i t s

Above some c e r t a i n f i l l e r

the

i n c l u d i n g m a t h e m a t i c a l model

o r i e n t a t i o n of anisometric

increase

lower

i s c l o s e l y related to

subsequent d e c r e a s e in t h e l a t e r a l dimension of v o i d s their

voids

c o n c e n t r a t i o n s r e l a t i v e l y e a r l y s a t u r a t i o n of void

v s draw r a t i o see

The amount o f

in the composite. For

and

compensating

low

filler

necking.

concentration

d e f o r m a t i o n mechanism c h a n g e s

[6],

( a b o u t 40 w t . % ) ,

saturation of the

the

relative

i n c r e a s e i n volume v s draw r a t i o o c c u r s much l a t e r o r even i s observed a t a l l

(see F i g . l ) .

On t h e m a c r o s c o p i c l e v e l t h e

d e f o r m s homogeneously. T h i s e f f e c t n e c k i n g on t h e m i c r o s c o p i c the f i l l e r

particles)

level

i s c a u s e d by y i e l d i n g and

( i n many p l a c e s a t o n c e ,

between

which makes d e c r e a s e i n t h e l a t e r a l

s i o n o f v o i d s much l e s s

dimen-

effective.

The a b o v e e x p e r i m e n t s p e r m i t t e d u n d e r s t a n d i n g o f t h e b a s i c r e s of the deformation process of f i l l e d

polypropylene.

indicated a l s o the p o s s i b i l i t y of o b t a i n i n g the highly polypropylene f i l l e d with modified c h a l k . (e.g.

such m a t e r i a l

on t h e t e c h n o l o g i c a l

semicrystalline

i s obtained

oriented

i.e.

in t h e c o n t i n u o u s

by d e f o r m a t i o n o f

process of f i l l e d

studied e x t e n s i v e l y using P l a s t i c i s e r

- the laboratojy set

[7J.

Fig.2

volume v e r s u s draw r a t i o

i l l u s t r a t e s the r e l a t i v e

experimental increase

(defined as a r a t i o of r o t a t i o n a l

o f t a p e s c o n t a i n i n g 40 wt.% o f m o d i f i e d

o r i e n t e d c o n t i n u o u s l y a t two t e m p e r a t u r e s : Relative

regulated

s p e e d s and t e m p e r a t u r e s . F o r d e t a i l s o f

of take-up r o l l s )

extruded

p o l y p r o p y l e n e has been

containing miniextruder with the take-up r o l l s of conditions see

process pure,

rolls.

Continuous o r i e n t a t i o n

rotational

of

reasons

l i n e used f o r o r i e n t a t i o n o f

thermoplasts),

s h e e t on t a k e - u p

Practical

featu-

They

m a t e r i a l w i t h h i g h amount o f v o i d s by s i m p l e d e f o r m a t i o n require that

not

sample

i n c r e a s e i n volume has

40°C and

measurements performed by w e i g h i n g i n a i r and i n d i s t i l l e d f o r not t o cause e r r o r

in d e n s i t y

chalk

100°C.

been c a l c u l a t e d from t h e

( a s w i l l be shown below w a t e r s o r p t i o n k i n e t i c s

in

speeds

density water

i s slow enough

determination).

145

I

S 13 fc I g

-

1.2 /

/

1.1

-

Qc

Fig.2.

h

x

/

/

/

2

y

/

o Td=M°C • Td =100° C I

3 4 5 DRAW RATIO

I

6.

I

7

Relative increase in volume vs. draw ratio during uniaxial orientation in continuous process. i-PP + + modified chalk 6:4, temperatures 40°C and 100°C.

Results shown in Fig.2

indicate that deformation mechanism at

40°C is rather similar to described above deformation mechanism of finite sample (INSTRON) at room temperature. Above 100°C the deformation mechanism changes which is connected with the a relaxation of polypropylene matrix occurring closely to this temperature and causing the change in the mechanism of yielding and deformation of matrix

[V] (in practice polypropylene is usual-

ly oriented at the temperatures close to the a relaxation which assures the best mechanical properties of an oriented material[8])• Fig.3 shows the morphology of the surface of an oriented tape observed in the scanning electron microscope. As it can be seen considerable amount of voids is opened at the film surface; in agreement with previous considerations anisometric chalk particles are oriented along the orientation direction. The presence of elongated voids enhances film fibrillation. Using rotating fibrillator interesting material is obtained which resembles in

146

Fig. 4.

SEM of single fibres of fibrillateci pure (left) and modified chalk filled i-PP.

147

touch and in appearance voids)

(white c o l o u r due t o l i g h t

the n a t u r a l plant f i b r e . F i g . 4

s c a t t e r i n g on

compares the morphology of

s i n g l e f i b r e s o f f i b r i l l a t e d pure and f i l l e d

i-PP.

Due t o the presence of v o i d s o r i e n t e d f i l l e d p o l y p r o p y l e n e

exhi-

b i t s s o r p t i o n of v a r i o u s l i q u i d s and vapours

see

[ 7 ] ) . Fig.5

illustrates

(for details

some r e s u l t s o f water uptake s t u d i e s

obtained f o r p o l y p r o p y l e n e c o n t a i n i n g 40 wt.% of m o d i f i e d chalk o r i e n t e d continuously a t 100°C (X=4, V/V 0 =1.15). R e s u l t s a r e plotted

in t h e form: f r a c t i o n of v o i d s f i l l e d w i t h water v s .

l o g a r i t h m of t i m e . Measurements were performed by continuous r e c o r d i n g of w e i g h t of

the sample suspended in d i s t i l l e d

water,

by means of INSTRON f o r c e measuring system. As i t can be seen, after

c o n s i d e r a b l e time,

the amount of absorbed water

indicates

t h a t a l l v o i d s take part i n the s o r p t i o n p r o c e s s . The r a t e of

-2.

0.

-/.

I

iog(t[h]) Fig.5.

Water u p t a k e ( f r a c t i o n of v o i d s f i l l e d w i t h water) in o r i e n t e d m o d i f i e d chalk f i l l e d chalk,

X =4).

i-PP.

(40 wt.% of

148 s o r p t i o n d e c r e a s e s w i t h time

( p e n e t r a t i o n depth) which i s a

consequence of v o i d system a n i s o t r o p y perpendicular t o the d i r e c t i o n of

( d i r e c t i o n of v o i d s

liquid front

is

penetration).

H y d r o p h i l i c chalk p l a y s an a c t i v e r o l e in water soption in hydrophobic p o l y p r o p y l e n e m a t r i x . The d r i v i n g f o r c e f o r c a p i l l a r y pressure- i s stronger

sorption-

in the channels formed by chalk

and p o l y p r o p y l e n e v o i d w a l l s than in t h e channels c o n s i s t i n g of p o l y p r o p y l e n e v o i d s o n l y . When the sample i s vacuum d r i e d a f t e r prolonged water s o r p t i o n and then again submitted t o water s o r p t i o n the i n i t i a l p e n e t r a t i o n r a t e drops down s i g n i f i c a n t l y (see F i g . 5

) . A p o s s i b l e e x p l a n a t i o n of

t h i s e f f e c t i s t h a t chalk

p a r t i c l e s a r e "washed o u t " from v o i d s l a y i n g c l o s e l y t o t h e f i l m s u r f a c e in the f i r s t mass of order of first

sorption process

2% i s observed a f t e r vacuum d r y i n g ) ,

and a t the

s t a g e s p e n e t r a t i o n occurs in the system c o n s i s t i n g mainly

of v o i d s w i t h p o l y p r o p y l e n e K i n e t i c s of nature of matrix

(decrease in the sample

liquid

[ 7 ] and,

(e.g.swells

effects,

walls.

l i q u i d s o r p t i o n depends a l s o on the physicochemical it),

if

liquid

i n t e r a c t s w i t h the polymer

may be accompanied w i t h

e . g . r e v e r s i b l e creep under load

interesting

[ 9 ] ) . These problems

lay

however beyond the scope of t h e p r e s e n t paper and w i l l be d e s c r i b e d in d e t a i l s

elsewhere.

Cellulosic

Fillers

in S e m i c r y s t a l l i n e

C e l l u l o s i c and l i g n o c e l l u l o s i c been i n v e s t i g a t e d

Thermoplastics

composites w i t h t h e r m o p l a s t i c s

in many l a b o r a t o r i e s but not so many s y s t e m a t i c

s t u d i e s were c a r r i e d out as in the case of mineral f i l l e d t e s . R e c e n t l y an i n t e r e s t i n g paper sing and

have

[10]

composi-

appeared in which p r o c e s -

p r o p e r t i e s of composites c o n s i s t i n g of

thermoplastics

and c e l l u l o s e - b a s e d f i l l e r s have been d i s c u s s e d . Very r i c h

litera-

ture concerning these composites i s quoted t h e r e . These systems a r e more complicated than mineral f i l l e r of high tendency f o r f i l l e r

c o n t a i n i n g composites because

aggregation.

Our r e s u l t s discussed here concern o n l y some p r o p e r t i e s of

filled

p o l y o l e f i n s with bleached s u l f a t e c e l l u l o s e f l o u r ( C F ) . Apparent 3 powder d e n s i t y of t h i s f i l l e r i s 0.34 g/cm , average aspect r a t i o

149 i s about 12. The d r y i n g of t h i s f i l l e r a t 105°C r e s u l t e d i n moist u r e content o f

0.2% b e f o r e composite p r e p a r a t i o n . The blends

were obtained in a small s i n g l e - s c r e w extruder a t 200°C. The highest f i l l e r

content obtained w i t h

Stress- strain

curves were recorded with INSTRON t e n s i l e

machine

( s t r a i n r a t e 5 mm/min,

polyolefins

was 50 wt.%. testing

oar-shaped samples 0.5 mm t h i c k ,

6 mm wide,20 mm l o n g , t e s t i n g a t RT and RH 50%). As former d i s c u s sion

concerned mainly the chalk f i l l e d

i - P P we r e p o r t here the

r e s u l t s f o r the i s o t a c t i c p o l y p r o p y l e n e Malen J-400. The d e n s i t y of the i - P P samples w i t h d i f f e r e n t f i l l e r c o n t e n t s was measured d e n s i t o m e t r i c a l l y and c a l c u l a t e d from weight d i v i d e d by volume. The experimental data are lower than the

theoretically

p r e d i c t e d . T h i s i s due t o i n e f f i c i e n t packing of f i l l e r

and t o

some gaseous i n c l u s i o n s . The d e n s i t y v a l u e s were in the range from i T 0.908 g/cm t o 1.08 g/cm f o r the samples with 0% t o 40% weight percent of the f i l l e r c o n t e n t

respectively.

Young's modulus dependence on f i l l e r content shows an i n c r e a s e from about 1 GPa t o 2.9 GPa by 50% of f i l l e r c o n t e n t . The v a r i a t i o n s of the s t r e s s a t y i e l d a r e not v e r y important w i t h the content up t o 20% and then, a t higher f i l l i n g

filler

(about 40-50%) o^

reaches 50% of the former v a l u e . Elongation a t break d e c r e a s e s w i t h the i n c r e a s e in f i l l e r content reaching 8% f o r samples w i t h 40% of CF. The dependence of the e l o n g a t i o n a t break i s w i t h the s i m i l a r decrease in the impact s t r e n g t h

drop in impact s t r e n g t h i s observed a l r e a d y a t r a t h e r lig

correlated

( I S ) . The sharp small

l e v e l but i t s d e c r e a s e per u n i t of f i l l e r added i s

smaller

f o r higher f i l l i n g d e g r e e . At the h i g h e s t l e v e l of f i l l i n g 2

IS i s of the order of

9 kJ/m . I t

seems i n t e r e s t i n g

i n f l u e n c e of moisture on IS i s small but the

the

t o note t h a t

treatment w i t h hot

water r e s u l t s in a more important d e c r e a s e in IS t o s u r f a c e d e f e c t s on the

fil-

Drobably due

surface.

Because of our i n t e r e s t in moisture uptake in chalk f i l l e d

oriented

i - P P we have i n v e s t i g a t e d more c a r e f u l l y t h i s e f f e c t in CF c o n t a ining systems. At RT and 50% RH the moisture uptake depends on filler

c o n t e n t . The s t o r a g e time f o r above mentioned

was about 100 days f o r reaching moisture content of sion of of

8% f o r

samples in hot water

conditions 1.7%. Immer-

(60°C) causes the moisture uptake

i - P P samples w i t h a f i l l e r

content of 40%. I t

b l e t h a t t h i s e f f e c t i s r e l a t e d t o the p l a t e - t h r o u g h

is

possi-

effect,i.e.

150 the d i s r u p t i o n of the continuous f i l m s u r f a c e . T h i s may be unacceptable in many a p p l i c a t i o n s but i t seems t o be of

i n t e r e s t when l o o k i n g f o r

strong e l a s t i c and porous m a t e r i a l s

obtained by s t r e t c h i n g , c a p a b l e of case the

plate-through

high water s o r p t i o n .

In t h i s

e f f e c t would be d e s i r a b l e .

Apart from much lower f i l l e r d e n s i t y the discussed system e x h i b i t s many s i m i l a r i t i e s w i t h the system c o n s i s t i n g of carbonate f i l l e r s .

Cellulose f i l l e r s exhibit

in

i - P P and calcium the matrix

r e l a t i v e l y high r i g i d i t y and s t r e n g t h . They do not enhance t h e degree of r e i n f o r c e m e n t because of

l e n g t h r e d u c t i o n of

f i b r e s broken by the shear f o r c e s during p r o c e s s i n g ,

longer

which

e x p l a i n s the s i m i l a r i t y between f i b r o u s and p a r t i c u l a t e fillers.

It

cellulosic

seems necessary t o p o i n t out t h a t the l a c k of

of c e l l u l o s i c f i l l e r s the p r o c e s s i n g

hardness

i s important from the v i e w p o i n t of wear of

machinery.

Comments on the R o l e of the F i l l e r Nature T y p i c a l aspect r a t i o of c e l l u l o s i c of chalk p a r t i c l e s .

It

fillers

(Mineral

vs.Cellulosic)

i s comparable t o

that

seems t h a t they could be used in the same

manner as e . g . chalk t o o b t a i n o r i e n t e d porous s o r p t i v e m a t e r i a l s , o f f e r i n g the a d d i t i o n a l advantage of the a p p r o p r i a t e i n t e r f a c i a l The concept of filler

t h e i r low d e n s i t y ,

provided

properties.

i n t r o d u c i n g a l i q u i d l a y e r between the polymer and

seems t o be u s e f u l a l s o in the case of c e l l u l o s i c

However the c o n d i t i o n s which have t o preserve

fillers.

be f u l f i l l e d in order

n e c e s s a r y i n t e r a c t i o n s between p o l y o l e f i n i c

to

matrices

and c e l l u l o s e - t y p e f i l l e r s are more c o m p l i c a t e d . The f i l l e r

sur-

f a c e should be w e l l wetted by the m o d i f i e r . The p o l a r nature of c e l l u l o s i c m a t e r i a l s imposes the use of not s w e l l the c e l l u l o s i c

polar l i q u i d .

It

f i b e r s or p a r t i c l e s because of

changes i n s t i f f n e s s . On the other hand the i n t e r f a c i a l

should possible agent

should p r o v i d e s u f f i c i e n t i n t e r a c t i o n s w i t h p o l y o l e f i n s .

Taking

these i n t o c o n s i d e r a t i o n one can e x p e c t t h a t o l i g o m e r s of

poly-

e t h y l e n e or p o l y p r o p y l e n e o x i d e s of a p p r o p r i a t e molecular

weights

are s u i t a b l e

i n t e r f a c i a l a g e n t s . They a r e l i q u i d s of

high b o i l i n g

p o i n t , thus t h e i r c o n c e n t r a t i o n on the s u r f a c e of f i l l e r s ought not change a t usual p r o c e s s i n g and s e r v i c e

temperatures.

151 The c o v e r i n g technique of c e l l u l o s i c

f i l l e r s w i t h above mentioned

o l i g o m e r s i s more d i f f i c u l t because one

cannot

use water

suspen-

s i o n , but a thinner of a low b o i l i n g p o i n t and s u f f i c i e n t high s o l u b i l i t y f o r epoxide o l i g o m e r s can be used. Drying a t t e temperatures should r e s u l t

appropria-

in w e l l covered m o d i f i e d CF.

The volume i n c r e a s e during the d e f o r m a t i o n of

such systems has not

been a l r e a d y determined e x a c t l y as a f u n c t i o n of CF c o n t e n t . From the d e n s i t y d e c r e a s e one can e s t i m a t e t h a t i t

i s of t h e order of

t h a t found f o r 1-PP f i l l e d w i t h 0E0 m o d i f i e d c h a l k . should c o n t a i n v o i d s ,

Such a system

nonuniformly o r i e n t e d polymeric matrix and

o r i e n t e d CF. Further s t u d i e s are needed in o r d e r t o compare the s t r u c t u r a l c h a r a c t e r i s t i c s of these systems and those obtained well

investigated

i - P P f i l l e d w i t h m o d i f i e d chalk o r i e n t e d

for

composi-

tes. I t has been shown t h a t voids are p a r t i a l l y

in the case of o r i e n t e d ,

interconnected

chalk f i l l e d

[ 7 ] . This f a c t e x p l a i n s

s o r p t i o n p r o p e r t i e s of t h e s e composites. Taking i n t o the p a r t i c u l a r

i-PP

high

consideration

s o r p t i o n a b i l i t y of c e l l u l o s i c m a t e r i a l s and t h e i r

s p e c i f i c i n t e r a c t i o n s w i t h p o l a r l i q u i d s a high f i l l e r

swelling

during s o r p t i o n can be e x p e c t e d . These p r o c e s s e s a r e connected w i t h a l a r g e v a r i e t y in

mechano-chemical e f f e c t s even in an

enhanced manner than those which a r e observed in the case of PP-chalk

systems.

References 1. Kowalewski, T . , A . G a l ^ s k i ,

submitted t o J . A p p i .

Polym.Sci.

2. G a l ^ s k i , A . , R . K a l i r i s k i . 1980. I n : Polymer Blends: P r o c e s s i n g , Morphology and P r o p e r t i e s v o l . 1 . ( E . M a r t u s c e l l i , M.Palumbo and M.Kryszewski, e d s . ) Plenum Press p.454. 3. K a l i i i s k i , R . , A . G a l Q s k i , M.Kryszewski. 26, 4047.

1981.

J.Appi.Polym.Sci.

4. Badran, B.M., A . G a l ^ s k i , M.Kryszewski. 27, 3669.

1982.

J.Appi.Polm.Sci.

5. Kowalewski, T . , R . K a l i r f s k i , A.GalQski, M.Kryszewski. C o l l o i d and P o l y m . S c i . 2 6 0 , 652.

1982.

6. Kowalewski, T . , A . G a l ^ s k i , M.Kryszewski. 1984. I n : Polymer Blends: P r o c e s s i n g , Morphology and P r o p e r t i e s v o l . 2 .

152

(M.Kryszewski, A . G a l ^ s k i , p.223. 7. Kowalewski, T . , A . G a l ^ s k i . 8. Samuels, R . J . , Wiley, p.85.

E.Martuscelli eds.).Plenum Press In p r e p a r a t i o n .

1974. I n : Structured Polymer

Properties,

9. Kowalewski, T . , A . G a l ^ s k i . 1985. Presented a t 17 th EPS Conference on Macromolecular P h y s i c s , Morphology of Polymers, Prague. 10. Kalson, C . , J.Kubat, H.-E. S t r f l m v a l l . 1984. 159.

J.Polym.Mat.10,

CELLULOSIC FILLERS FOR THERMOPLASTICS

C. Klason, J. Kubat Department of Polymeric Materials, Chalmers University of Technoloqy S-412 96 Gothenburg, Sweden

Introduction Cellulosic and lignocellulosic materials show interesting potential as fillers for thermoplastics. Although their use in thermosets is a well established technique, the area of thermoplastics-based composites seems to have attracted relatively little attention by processors and compounders. During recent years, however, there appears to be growing interest in this area, at least judging from the increasing amount of relevant literature. Among the properties of such fillers and reinforcing agents, which may be of interest in the present context, one can mention their relatively low den3 sity (c. 1500 kg/m ) and reduced wear on the processing equipment. In addition there is the possibility of disintegrating the fibre into microscopic fibrils, after proper pre-treatment (hydrolysis), which fibrils can show stiffness and strength characteristics in the vicinity of those found for carbon fibres. The present paper summarizes the main results of a systematic study of processing and mechanical parameters of some typical composites based on a number of common thermoplastics and various cellulosic fillers and lignin. Special attention is given to the possibility of utilizing the embrittlement imparted to native cellulose by a hydrolytic pre-treatment (chain length reduction) to improve homogeneity and also the mechanical parameters of the composites produced. Indeed, the findings reported below support the expectation that the potential of cellulose fillers may be significantly improved by such pre-treatment. The thermoplastics used were low and high density polyethylene (LDPE, HDPE), polypropylene (PP), polystyrene (PS), Doly(methyl

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

154 methacrylate)

(PMMA), and polyamide 6 (PA6). The fillers were

bleached cellulose fibres and flour, wood flour, lignin, pre-hydrolyzed cellulose fibre, and a commerical microcrystalline cellulose (Avicel). Since such fillers are normally not compatible with the plastics listed above, a number of coupling agents were tried out, a PP-maleic acid copolymer appearinq to be especially efficient. Some experiments were also conducted with grafted cellulose fibres in a PMMA-matrix. Relevant literature data have been reported in connection with earlier publications on this subject

(1,2,3,4,5).

Experimental The fillers used were cellulose fibres of bleached pine sulphate, bleached pine sulphite, bleached hardwood

(birch and beech) sul-

phite, bleached dissolving pulp (pine sulphite, alpha cellulose, MoDo Crown), cellulose flour (bleached birch sulphate), wood flour (spruce) and Kraft lignin (Indulin AT, Westvaco). The cellulose fibres, supplied as pulp bales, have to be qround in a strand pelletiz er (Rapid type GK 20) in order to facilitate the feeding of the fibre-polymer blend into the extruder. The average fibre length after grinding is about 1 mm (aspect ratio 40); for the other filler the average length (ym): aspect ratio was 100:2.5 (wood flour), 100:4.3

(cellulose flour) and 4:1 (lignin, spheri-

cal) . The fillers have to be carefully dried

(105°C for 12 h, or

80°C in vacuum for 24 h) before compounding, otherwise a pronounced discolouration of the material occurred. Also, the melt temperature had to be kept as low as possible, preferably below 200°C. The polymers used were LDPE (DFDS 4401, Unifos Kemi, density 918 kg/m 3 , MI,-value 2), HDPE (DMDS 7006, Unifos Kemi, density 960 3 kg/m , MI 2 -value 7), PP copolymer (GYM 621 powder, ICI, density 905 kg/m 3 , MI 2 ~value 13 at 230°C), PP homopolymer

(Profax PC 072,

3

Hercules, density 905 kg/m , MI 2 -value 3.5 at 230°C), PS (N 4001, Hoechst, density 1050 kg/m 3 , MIg-value 4.5 at 200°C), PMMA (N 7, Rohm, density 1.17 kg/m 3 , Mi-value 6 at 230°C), and

155

PA 6 (Grilon A 28 GM, Emser, density 1130 kg/m 3 , T^ 215°C). PMMA and PA 6 were dried in vacuum (80°C) for 24 h before compounding. In some experiments, the cellulose fibres were hydrolyzed in solution (5% H-^SO^, 100°C for 1 h) , neutralized with ammonia, and washed and dried (50°C for 48 h, and 105°C for 12 h). The hydrolytic treatment reduced the DP value of the cellulose from about 1400 to about 250. The resulting embrittlement facilitated the disintegration of the fibres into smaller fragments in the shear field of the compounding and processing machinery, and also improved the degree of their dispersion in the matrix. For comparison, a microcrystalline cellulose powder (Avicel PH-101, FMC, particle size about 50 pm) was compounded with HDPE and PP. For some experiments, bleached pine sulphite fibres (MoDo Crown) were modified by grafting with methyl methacrylate (MMA) , grafting level 20%, to make them compatible with a PMMA-matrix. The grafting was made with redox initiator using Mn ion complexes in aqueous solution at room temperature (RT) (6). The components were homogenized in a compounding extruder (Buss Kneader PR 46 in experiments with unmodified fillers, Werner & Pfleiderer ZSK 30 twin-screw extruder in experiments with hydrolyzed fibres). The extrudate strand was granulated prior to the injection moulding step (Arburg 221E/17R). The test bars (DIN 53455) with a cross-section of 10 x 3.5 mm and an effective length of 75 mm were conditioned at RT and 50% rel. humidity for 24 h before tensile testing (ASTM D 638, Instron Model 1193) at RT and 4.5 >10 3 s deformation rate. From the stress-strain curves, the tangent modulus E, the strength at yield and rupture, a^ and a^, and the corresponding deformations, e^ and e^, were determined. The impact strength was measured with a Zwick instrument (Model 565 K, unnotched and notched Charpy, DIN 53453). Filling levels are given as weight percentages.

156

Results Figs. 1-3 reproduce the main results of the experiments with unmodified cellulosic fillers, as reported to some extent in refs. (1,2,3,4,5) .

Figure 1. Tensile modulus and yield strength vs. filler content for HDPE. Compounding in Buss Kneader except for cellulose fibre (bleached pine sulphate in ZSK). Fig. 1 shows the effect of the various fillers when incorporated in a HDPE-matrix. As can be seen, cellulose fibres produce the highest modulus increase, followed by wood and cellulose flour, microcrystalline cellulose and, finally, Kraft lignin. The high efficiency of cellulose fibre appears to be due to their high aspect ratio. It may be noted that in this case the compounding was carried out with the ZSK-machine, the single screw compounding extruder (Buss) being unable to produce a satisfactory dispersion. The twin-screw compounder, equipped with several mixing elements along the length of the screw, reduced the fibre length from 1.0 to 0.5 mm, producing a good dispersion. On the other hand, for the ground fillers (flour products) the dispersion efficiency of the Buss and ZSK extruders was about equal. The low stiffening effect obtained using lignin appears to be related to the low modulus value of this material (6.6 GPa according

157

to refs. (4) and (7). Among the other particulate fillers, wood flour stands out as the material with the highest stiffening capacity. It may be fair to assume that the licmin content of this material has a positive influence on adhesion to the matrix. The right-hand part of Fig. 1 shows the change in the yield stress with the filler content. As can be seen, the materials with the highest stiffening effect (cellulose fibre, wood flour) produce an increase in a , while the opposite is true of microcrystalline cellulose and Kraft lignin. Cellulose flour takes an intermediate position in that it leaves a unchanged up to more than 40% filling level.

0

20

40

o FILLER CONTENT,

20

40

60

%

Figure 2. Elongation at break and impact strength (unnotched, Charpy) vs. filler content for HDPE. Symbols as in Fig. 1. As commonly observed with filled systems, the impact strength and the ductility of the samples discussed in connection with Fig. 1 were negatively affected by the fillers. Fig. 2. However, there are two exceptions. The high aspect ratio of the cellulose fibres gives a slightly better ductility and, secondly, the low modulus of lignin appears to result in a plasticizing effect (below 30% filler content).

158

FILLER

CONTENT,

*

Figure 3. Tensile modulus vs. wood flour content for different thermoplastics. Compounding in Buss Kneader. Fig.3 summarizes the stiffening effect of wood ous matrix materials used. As can be seen, the tent curves are nearly parallel, implying that ening effect decreases with increasing E-value

flour on the varimodulus-filler conthe relative stiffof the matrix.

Figure 4. SEM micrograph of a fracture surface of PP/wood flourcomposite without (left) and with (right) 6% coupling agent (maleic anhydride modified PP). Magnification 300. Length of white line 100 pm.

159 A number of additives have been tested for their efficiency as coupling agents improving the filler-matrix adhesion (2). A PPmaleic acid copolymer (8) was found to be particularly suitable for this purpose. Fig. 4 illustrates the difference between treated and untreated wood flour in a PP-matrix. The pull-out effect observed in a fracture surface is significantly reduced by adding 6% of the coupling agent to the filler.

Figure 5. Tensile modulus, stress and elongation at break for dry and conditioned PA6 containing cellulose fibres. Fibre content 40%. Conditioning: 80°C in water for 20 h. Compounding in ZSK. • PA6; filler: • bleached pine sulphite, A bleached pine sulphate. Unfilled symbols - unhydrolyzed, filled symbols - hydrolyzed. Polyamides can be expected to qive a fair degree of adhesion to cellulosic fillers, since hydrogen bonds represent the major bonding mechanism in both cases. Fig. 5 confirms this for different cellulose fibres (40% filling content) in PA6. The hygroscopicity of both components of such composites results in the well-known reduction in E and increase in ductility and impact strength when such samples are conditioned (80°C in water for 20 h), Figs. 5 and 6. These figures also indicate that pre-hydrolyzed cellulose absorbs less moisture than the untreated material.

160

Figure 6. Impact strength (notched, Charpy) for the samples of Fig. 5 (same notations). An obvious means of improving the filler-matrix adhesion is by grafting the molecules of the matrix material onto the filler particles. This has been reported earlier for bagasse (9). Also impregnating the filler (sawdust) with the corresponding monomer followed by polymerisation (10) had a beneficial effect on the adhesion. However, injection moulding of such composites has not been reported. In our experiments, PMMA was compounded with up to 25% MMA-grafted cellulose fibres and injection moulded into test bars. Although this material combination hardly can be considered a candidate for commercial applications, the results obtained show that grafting is a highly efficient method for improvincr the filler-matrix adhesion. This is reflected in an increase in the tensile strength values. Fig. 7, and also in a decrease in fibre pullout as evident from the SEM micrograph of a fracture surface, Fig. 8.

As already mentioned, pre-hydrolytic treatment of the cellulosic filler offers a number of interesting possibilities with regard to improving the mechanical parameters and the degree of dispersion of the corresponding composites. The well-documented embrittlement (11) of the cellulose material due to such pre-treatment can be expected to result in a significant disintegration of the fibres or particles by the shear forces acting in normal processing equipment .

161

PERCENT in PMMA

GRAFTED

CELLULOSE

Figure 7. Tensile modulus, stress and elongation at rupture, and impact strength (notched, Charpy) vs. filler content for PMMA containing grafted cellulose (20% MMA grafted on the bleached pine sulphite fibre). Compounding in ZSK.

Figure 8. SEM micrograph of a fracture of a PMMA sample filled with 12.5% grafted cellulose. Magnification 300. Length of white line 100 ym.

162

6

-

5

-

Bh c„

. . . ...

4

(3 -3 UJ

A

1 n

\ \

i of 0.

•

O

2

B

Ou

Q.

1

\ \ > \

\ \ \ \

\ 1

1

1

1

1

1

V \

Figure 9. Tensile modulus for PP (copolymer 1, homopolymer 2) without and with 40% cellulose fibres, h - hydrolyzed material (bleached fibres). A - pine sulphite, B - pine sulphate, C - hardwood (birch and beech) sulphite, D - alpha cellulose, E - alpha cellulose in PP homopolymer with 3% added coupling aqent (maleic anhydride-PP copolymer). Compounding in ZSK. Fig. 9 shows that such an improvement is, in fact, obtained, in this case with a PP-matrix. At the 40% fillincr level, modulus values in the vicinity of 5 GPa are obtained with pre-hydrolyzed fibres compounded with a PP copolymer. These values can be increased somewhat when using PP homopolymer together with a coupling aqent (PP-MA-copolymer). Microscopic investigations showed that the main disintegration mechanism of the fibres was irregular fragmentation combined with a certain amount of formation of fibrillar entities with varying length and aspect ratio.

Figure 10. Strength at yield for samples in Fig. 9 (same notations).

163

The beneficial effect of the pre-hydrolysis on the stiffness of the composites was not paralleled by their ductility and impact strength, both these parameters being reduced as demonstrated in Figs. 10-12.

Figure 11. Elongation at yield for samples in Fig. 9 (same notations) . The slight differences in the mechanical parameters related to the origin and morphology of the fibres used were almost entirely eliminated by the hydrolytic treatment. The mechanical degradation of the PP-matrix material due to compounding and injection moulding was negligible (GPC measurements).

20

to 10

5

o Figure 12. Impact strength (notched, Charpy) for samples in Fig. 9 (same notations).

164

Final Remarks This brief account of the mechanical parameters of some typical compounds based on common thermoplastics and cellulosic fillers shows but a few features of this group of composite materials. On the whole, particulate cellulosic fillers appear to act in the same way as the corresponding mineral materials used for this purpose as, for instance, calcium carbonate. When the length of the cellulose fibre is retained during processing (twin-screw compounding) , the modulus is significantly improved, althoucrh a reduction of ductility and impact strength must be taken into account. However, such a balance between modulus (stiffness) and ductility applies only to unmodified fillers. By providing a suitable interface between the filler particles and the matrix, for example by using coupling agents or rubbery layers, substantial improvements in the modulus-ductility balance can be expected. In the present case, only some preliminary tests have been made. The main result of this report is the demonstration that a hydrolytic treatment of the cellulose fibre and the resulting embrittlement produces significant improvement in the modulus (PP matrix). The E-values recorded for such composites exceed the values obtained with carefully compounded, untreated fibres. Although the main part of the embrittled material appears to occur in the matrix in the form of irregular fragments, a disintegration into high-modulus, high-strength fibrils is discernible. Control of the conditions under which such fibrils are formed from the pre-hydrolyzed material offers interesting possibilities for using the inherent properties of these building blocks of the native fibre to achieve significant reinforcement effects in composites. Calculations using the Tsai-Halpin formula (12) show that modulus values in the vicinity of 10 GPa can be expected for PP-based materials when the cellulose phase is converted into such fibrils. This prediction is supported by the fair agreement between the predictions of the Tsai-Halpin theory and the present results, especially when the distribution of the aspect ratio values has been taken into account.

165 Acknowledgement Financial support from the National Swedish Board for Technical Development is gratefully acknowledged. Thanks are also due to Professor B. Ränby, Royal Institute of Technology, for preparing the grafted cellulose fibres, and to Mr. A. Boldizar for skillful experimental assistance.

References 1. Klason, C., J. KubSt, H.-E. Strömvall. 1984. Int.J.Polym.Mat. 10, 159. 2. Dalväg, H., C. Klason, H.-E. Strömvall. 1985. Int.J.Polym.Mat. 11, 9. 3. KubSt, J., H.-E. Strömvall. 1983. Plast.Rubber:Proc.Appl. 3, 111.

4. Klason, C., J. Kubät. Submitted to Plast.Rubber:Proc.Appl. 5. Berggren, K., C. Klason, J. KubSt. 1975. Kunststoffe 12, 69. 6. Ränby, B., L. Gädda. 1982. In: Graft Copolymerization of Lignocellulosic Fibers. ACS Symposium Series, No. 187 (D.N.S. Hon, Ed.). Washington, pp. 33-43. 7. Cousins, W.J. 1977. N.Z.J.For.Sei. 7:1, 107. 8. Brit. 1,101,408. 9. Nagaty, A. 1979. J.Appl.Polym.Sei. 23,

3263.

10.U.S. 3,083,118. 11.Hägglund, E. 1939. Holzchemie. Akademische Verlagsgesellschaft M.B.H. Leipzig, pp. 74-96. 12.Lewis, T.B., L.E. Nielsen. 1970. J.Appl.Polym.Sei. 1£, 1449.

THE EFFECT OF FILLERS ON THE RHEOLOGICAL AND MECHANICAL PROPERTIES OF POLYPROPYLENE COMPOSITES

B. Puk£nszky, F. TiidSs, T. Kelen Central Research Institute for Chemistry, Hungarian Academy of Sciences, H-1525 Budapest, P.O.Box 17, Hungary

Introduction Polymer composites containing different type and amount of fillers have great importance, both theoretically and practically. The majority of manufactured thermoplastics is also available in filled or reinforced grades. Polypropylene (PP) is not an exception and the interest in filled PP is reflected by the increasing number of publications in this field. Like the other thermoplastics, PP can be filled with mica (1-5), CaC03 (6,7), silica (8) and other fillers (9) . The characteristics of filled polymers differ significantly from those of unfilled materials. The changes in properties are determined, among other things, by the size, shape, chemical composition, surface, etc., of the filler. Although considerable information is available on the effect of these factors, there are still numerous unsolved problems. The aim of the present work was to investigate the effect of different fillers on the properties of PP composites and to determine the most important factors which influence the mechanical and rheological characteristics of filled plastics.

Experimental In the work Tipplen H501 (TVK, Hungary) was used as polymer. As fillers, five CaCO^ products from different sources, a mixture of chalks to obtain a broad particle size distribution, three talcs, a mica and a silica were investigated. The composites contained 0.2 % Irganox 1010 (Ciba-Geigy, Switzerland) antioxidant. The investigated fillers and their most important characteristics are listed in Table 1.

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • N e w York - Printed in Germany

168 T a b l e 1. C h a r a c t e r i s t i c s of t h e S t u d i e d Trade name

Manufacturer

Socal U1

Solvay

Polcarb Calcilit 8

English Clays

Millicarb

Alpha Ctnya

C a l c i l i t 100

Alpha

Mixture 3 Finntalc M05 Finntalc M15 Finntalc PF

a

Finnminerals Finnminerals Finnminerals

F i l l e r type CaC03 II fl II II talc II II b mica silica'3

Fillers

Density p 3 (g/cm ) 2.65

Particle s i z e Oil No. a t 50 w% Fo (g/100 g) (pm) 0.08 50.2

2.64 2.71 2.65

0.9 7.2

2.71

78.0

12.8

2.68 2.85

0.6 2.0 5.1

30.4 56.2

2.80 2.77 2.77 2.05

4.0

12.5 75.0 5.6

22.5 22.0 21.4

40.0 29.4 65.3 59.5

Socal U1/Polcarb/Calcilit 4: 33/33/34 v%.

^ Mica and fused s i l i c a were supplied by the State Research I n s t i t u t e of Materia l s Prague, Czechoslovakia.

C o m p o s i t e s were p r e p a r e d i n t h e Rheomix 600 m i x i n g chamber o f a HAAKE Rheocord EU-10V (HAAKE I n c . , USA) p l a s t o g r a p h . The m i x i n g c o n d i t i o n s were 190 ° C , 40 r . p . m . , 45 ml c h a r g e volume and 15 min. Comp o s i t e s c o n t a i n i n g 0 . 0 5 , 0 . 1 , 0 . 2 , 0 . 3 , 0 . 3 5 , 0 . 4 and 0 . 4 5 volume f r a c t i o n o f f i l l e r were p r e p a r e d i n most c a s e s . During t h e m i x i n g t o r q u e (M) and m a t e r i a l t e m p e r a t u r e (T) were r e c o r d e d a s a f u n c t i o n of t i m e . Torque a t 185 °C (M-jgg) p r o p o r t i o n a l t o v i s c o s i t y and s l o p e of In M v s 1/T c u r v e s ( f l o w a c t i v a t i o n e n e r g y , E^) were u s e d t o c h a r a c t e r i z e the r h e o l o g i c a l p r o p e r t i e s of the c o m p o s i t e s . C o m p r e s s i o n molded p l a t e s were p r e p a r e d a t 190 ° C ; s p e c i m e n s were c u t from them f o r t e n s i l e t e s t i n g . T e n s i l e t e s t s were c a r r i e d o u t a t 5 mm/min c r o s s - h e a d s p e e d . For i m p a c t t e s t i n g n o t c h e d Charpy b a r s were machined o u t of 4 mm t h i c k c o m p r e s s i o n molded p l a t e s . DSC measurements were c a r r i e d o u t w i t h a DuPont 910 (DuPont, USA) DSC c e l l w i t h 20 ° C / m i n h e a t i n g and c o o l i n g r a t e s , r e s p e c t i v e l y , t o d e t e r m i n e t h e m e l t i n g and c r y s t a l l i z a t i o n c h a r a c t e r i s t i c s of t h e composites. O i l number of t h e f i l l e r s u s e d t o c h a r a c t e r i z e t h e s p e c i f i c was d e t e r m i n e d a c c o r d i n g t o DIN 53 199.

surface

169

Results and Discussion

The rheoloqy of suspensions and filled polymers were widely investigated in many laboratories, but the rheological characteristics of composites containing more than a few percent of fillers cannot be described by simple mathematical equations. These systems show a non-Newtonian behaviour, time dependent rheological properties and yield stress at low shear (10). Most of these characteristics result from the agglomeration tendency of the filler (11). In Fig.1 the M^g^ torque values and in Fig.2 the flow activation energies are plotted for PP composites containing five CaCO^ fillers of different particle size. The torque value proportional to viscosity increases significantly with decreasing particle size. This observation is in accord with the results of White (10,12). The same tendency can be observed in the case of the flow activation energy. These latter results, however, must be treated with caution. The flow activation energies determined from the In M vs 1/T curves agree well with the activation energies measured by other rheological measuring devices, e.g., ca-

0

Fig.1.

0.2

O.U

fp

torque values of CaCO3 containing PP composites as a function of the filler volume fraction. Average particle size: x: 0.08 um; o: 0.9 ym; A: 4.0 ym; +: 7.4 um; V: 78 pm.

170

Fig.2. Flow activation energy of CaCC^ containing PP composites as a function of the filler volume fraction. Symbols as in Fig.1.

pillary viscometers, rotational viscometers, etc., for composites with low filler content. But the above mentioned agglomeration tendency of the filler (11) and the time dependence of the rheological characteristics make the determination of the flow activation energy difficult and very inaccurate. The torque values (^35) of the talc-filled composites also increase with increasing filler content, but no difference was observed between the effects of talcs with different particle size. This indicates that talc which has an anisotropic particle shape behaves differently from CaCO^. The limited results on the rheological properties of PP composites containing different fillers show that particle size, shape, and filler type all have a significant effect on these characteristics. Young's modulus of the same composites plotted in Fig.3. also increases with the filler content and with decreasing particle size. Numerous attempts have been made to describe the physical properties of polymer composites as a function of filler content. For our treatment we grouped these models according to their capability of describing the observed size effect of the filler.

171

Fig.3. Young's modulus of CaCO3 containing PP composites as a function of the filler volume fraction. mixture of chalks, average particle size: ~0.6 ym. Further symbols as in Fig.1.

Those models which contain only the physical characteristics (modulus, Poisson ratio) of the two components, naturally, cannot account for such effect. The most often used formula of this type is the Kerner's equation (13): 1 + ABcp f

G

G^ where A

=

7 - 5v = 8-10 v m

(1)

1-B o c a>

cr

*10

I I

0.8

0.9

1.0 Fi ber

1.1 2.6 length

2.8 3.0 ( mm )

tL

3.2

Fig.l. Frequency distribution of discontinuous glass fibers.

Table 1.

Reinforcing Fiber Length

Sample number

I

IE

HI

^

Mean fiber length (L)mm

0.90

2.06

2.94

5.0 3

Mean aspect ratio (L/d)

71

162

2 31

395

This procedure enables melting of polypropylene fibers for matrix materials, and obtaining preparation of bubble-free resins reinforced with random-planar orientation of discontinuous fibers. In this experiment, the volume fraction of the glass fibers was set at 11.4 %. The test specimens which measure 1.0 mm in thickness, 15 mm in width and 120 mm in length were cut from these composites and subjected to tensile test at a strain rate of 0.2/min with the aid of a Tensilon UTM-I-2500 (Toyo Baldwin).

Twelve specimens were used

for each fiber length and each temperature level tested.

186 Results and D i s c u s s i o n The r e l a t i o n s h i p between t e n s i l e s t r e n g t h and a s p e c t r a t i o f o r the p o l y p r o p y l e n e - d i s c o n t i n u o u s

glass

i n F i g u r e 2.

the t e n s i l e

At both t e m p e r a t u r e s ,

r a p i d l y as the a s p e c t r a t i o i n c r e a s e s

strength

increases

(2,10),

t h a t the t e n s i l e

i n which discontinuous

fibers

t a k i n g the thermal

strength

[aes]T

of

f a c i a l r e g i o n and i s w r i t t e n as

Fig.2.

depends inter-

:

(1~lc / 2 L ) g f

^j2+|n|T(L/d)a.v,(%)

12

1

23 "C

w

04

vims ) '83KT4 o&S-IOr* 0 83-1CT6 a8-3-10"7

\ \

b^tB^yH-B^ .

0-2 _ 0 0

Fig.7. At elevated temperatures the W , corresy/c ponds to $ m .

A0*C

-1

«fct, 08 - % 06

«a

ill-Bit lw-2^ 3

\\ N\

serie IIA...IIH \ 02

04

06 §f

60-0

; \; \ . serie IA..JH

i

3 y/mflTr^m and temperature. Mixtures P to V (e.g.), when investigated by measuring complex modulus and creep, exhibit no effect on E when the filler content is low (where

°y/ c /

a

y/m

>

'

^ a m Pi n i 3 i s

a

little higher

than that of the matrix, as well as the low stress creep. A low filler concentration induces probably some small strain nonlinearity. Reasons for the different behaviour of low filler content composites could be found well below the yield point. If, due to

214 sufficient interparticle distance, the coalescence of "around the filler-voids" is less probable (10,11), the energy consuming processes can act in the matrix over a broader range of strain. Moreover, at higher total strains, activity of the particles as barriers of the local yield is to be expected.

Conclusions The general conclusion that the poor

filler/matrix adhesion is an

important reason for the yield stress drop due to higher contents of the particulate filler, is supported by this work. On the other hand, considering the low strain creep at elevated temperature, the advantage of a higher filler content should be emphasized. An important factor in the optimization of polypropylene composites is the matrix itself. Considering the "tensile strength", polymers with sufficient ductility seem to be more tolerant to non-reinforcing fillers. Because of the rather complex interactions of components, the development of thermoplastic composites is affected by the great number of experiments required. To reduce the amount of testing, an equation has been proposed which contains easily determinable variables. A secondary use of Eq.3 can be recommended - to check the filler/matrix adhesion through the determined value of B. In the course of developmental work, natural deviations from Eq.3 should be distinguished from the consequences of deliberate changes in the composition. However, further questions are still open for research.

References 1. Ishai, 0., L.J. Cohen. 1968. J.Composite Materials .2, 302 2. Evans, A.G. . 1972. Phil.Mag. .26, 1327. 3. Sultan, J.N., F.J. McGarry. 1973. Polym. Eng. Sci. 13., 29. 4. Mallick, P.K., L.J. Broutman. 1975. Mater.Sci.Eng. 18, 63. 5. Landon, G., G. Lewis, G.F. Boden. 1977. J.Mater.Sci. L2, 1605. 6. Young, R.J. , P.W.R. Beaumont. 1977. J.Mater.Sci. 12., 684. 7. Green, D.J., P.S. Nicholson, J.D. Embury. 1979. J.Mater.Sci. 14, 1413.

215 8. Idem, ibid, p. 1657. 9. Theocaris, P.S., G.C. Papanicolaou, G.A. Papadopoulos. 1981. J.Composite Materials lj^, 41. 10. Friedrich, K., U.A. Karsch. 1981. J.Mater.Sei. 16^ 2167. 11. Idem. 1982. Polymer Composites _3, 65. 12. Kuiera, J.. 1983. Plasty a Kauiuk 20,

289.

13. Moloney, A.C., H.H. Kausch, H.R. Stieger. 1983. J.Mater.Sei. 18, 208. 14. Ramsteiner, F., R. Theysohn. 1984. Composites ljj, 121. 15. Spanoudakis, J., R.J. Young. 1984. J.Mater.Sei. 19, 473. 16. Maxwell, D., R.J. Young, A.J. Kinloch. 1984. J.Mater.Sei.Letters 2, 9. 17. Charrier, J.-M.. 1975. Polym. Eng. Sei. 15^, 731. 18. Piggot, M.R., J. Leidner. 1974. J.Appl.Polymer Sei. 18_, 1619.

THERMOELASTIC EFFECT OF "POLYPROPYLENE - CaCOj" COMPOSITES. THE INFLUENCE OF THE COMPOSITION, RATE OF STRAIN AND TEMPERATURE

J.Hugo, M. Houskovi, V. Matëna National Research Institute for Materials 113 12 Praha 1, Opletalova 25

Introduction Some aspects of the

strength and toughness of polymer composites

are associated with the

energy consumption in relatively harmless

processes during loading. These study of the

factors may be important in

interactions of polymeric

the

matrices and dispersed

hard or elastomeric particles. The energy is consumed primarily by the polymeric matrix

(1), and the most

important mechanisms

are yielding and crazing. In small deformations (2), the total strain tensor, e , can be decomposed into thermodynamically reverd e sible deformation, e , deformation e resulting from creation of discontinuities and

e 1 , irreversible deformation. The total work

of deformation, W, thus has three corresponding components

W

=

W

e

+

W, d

+

W. l

(1)

Part of VT is dissipated in damage formation and the rest of W^ is converted into heat. The viscoelastic behaviour of the polymeric matrix and the complexity of the composite structure are responsible for the superposition of several processes manifested in smooth continuous dependences of mechanical quantities. Thermal effects accompanying the deformation are relatively easily accessible for direct measurements. While uniaxial elastic elongation is associated with a temperature decrease (provided the coefficient of thermal expansion is positive) , the temperature of material deformed plastically under adiabatic conditions rises (3). For metals, an adiabatic change of temperature in the elastic range is described by the equation of W. Thomson

(1851)

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

218 a T Aas T

* s = -

' C

where

a

(2)

a

is the coefficient of linear thermal expansion, T the ab-

solute temperature,

Aa g is the isentropic change in stress and C a

the heat capacity per unit volume at constant stress. A thermoelastic temperature decrease results from a Poissons volume increase in a

uniaxially deformed body. There is much evidence

that, after the volume of the polymer increases to the critical value, conformational changes can begin and, consequently, crazes or yield appear. Thus, the requirement for conformational changes is the attainment of a critical volume, V (1 + a T ) (4). Semio g g crystalline polymers very probably need a necessary amount of total accumulated energy ( AH^) to overcome van der Waals bonds in the crystalline regions to allow conformational changes (4). Conditions in the micro-volume are somewhat complicated when a filler is added. Checking the toughness of thermoplastic composites by computing the energy consumption at strains far behind the yield point may distort information on the events between the "elastic limit" and the yield point, where the "damage" originates.

Experimental Procedure Materials Calcium carbonate powder Durcal 2 (OMYA) was used as a filler in all compositions tested. Isotactic polypropylene MOSTEN (melt flow rate 2.5 g/ 10 min ..21 N, 230

C) was a basic matrix for materials

P...IB. Materials AB, AC and AD contained polypropylene of reduced mol. weight (of low melt viscosity) with 28% by weight of EPDM masterbatch. Weight percent of CaCC>3 were: P...0/ U...21.5/ IY...9.5/ X...42.2/ IA...12.5/ IB..,26.5/AB...0/ AC...20/ AD...30. Materials IA and IB contained 24.4 % (by weight), or 15 % resp., of cut glass fibres of 13 urn diametre. The length of fibres after the specimen preparation was 0. 3 - 0. 4 mm ( = 1 mm). No special treatment was used on

219 the fibres. All composites were prepared by vtfMCH in Brno and by VUGPT in Gottwaldov. An optimized regime of injection molding was used for the preparation of test specimens of the ISO 527 type. A four to six months standard conditioning period was maintained before testing. Apparatus and method The cylindrical thermostatic chamber (Fig.l) is mounted on the INSTRON tensile tester. The temperature inside the chamber is kept constant within i 20 mK by gaseous N 2 flowing through the special stainless grips. The temperatures of the tested (T) and reference (T ) specimens are measured by the thermocouple battery couples are embedded into the specimen surface) as

(thermo-

AT (= T - T ).

AT in the unloaded state is less than 1 mK. The maximum

AT varies

according to the material tested and the testing conditions from 400 mK to 600 mK. The thermocouple signal is amplified by the Keithley 148 nanovoltmetre and recorded simultaneously with loading force F.

Fig.l. An apparatus for measuring the temperature changes of the loaded specimen.

220 Results and Discussion Only a few metals obey Eq.2 up to their yield point where the temperature change undergoes a sharp reversal (3). Considering the structure of polymer composites, processes

could be expected in

the range of decreasing temperatures, which dissipate part of the supplied energy, with some conversion into heat. The initial characteristic drop in temperature will gradually decrease with increasing elongation

A1 up to a minimum temperature which is the point

of temperature change reversal. Function "d(AT)/d(Al) versus F" is computed from the dependence of AT

and

A1 on F;

d(AT)/(.il) is denoted by the symbol dZ.

Several stages may appear on the computed "thermomechanical" curve: dZ is independent of F, dZ increases with increasing F and dZ decreases with increasing F. The increase of dZ is sharp or gradual. If the negative value is independent of F, a reversible deformation is assumed. An increase in dZ with increasing F indicates some energy dissipation. For dZ = O (excl. d(Al)

-»-

, d (AT)

= 0. The

cooling and heating processes are in equilibrium. The increase in dZ in the positive range is rapid, and F reaches its maximum. If d.Z decreases eventually in the negative range, some hardening is to be expected from the structural changes. Because of a certain degree of anisotropy, a simple continuous change in dZ is not very probable. Local yielding, voids, craze or crack nucleation, all of statistical character, may influence the dZ value. Moreover, a limited stability of the measuring system as well as artefacts of the computation contribute to some uncertainity in the interpretation. Hence, e.g., independence of dZ on F may sometimes be apparent, as a result of partial local yielding and hardening. The following conclusions should thus be considered as preliminary. Fig.2 depicts the dependence of dZ vs. F for unfilled polypropylene. At -20 °C and 0 °C the reversal to yield is sharp, appearing at a relatively high stress. At 23 °C and 40 °C, dZ increases

grad-

ually from nearly the commencement of loading. No qualitative difference can be expected in the mechanism of both cases. It seems that they differ in the volumes in which the dissipative processes are active and in the energy available to the stressed body. Simplifying the problem, a gradual increase in dZ is associated with conformational deformation of the micro-volume and with the

(par-

221

Fig.2. Thermomechanical curves dZ vs. F for the unfilled polypropylene.

Fig.3. Effect of the rate of strain, PP-CaCO^-

allel) cummulation of defect nuclei of a low total amount of yielding. A sharp rise in dZ should then be associated with a rapid sequence of the previously mentioned events, followed by spontaneous yield of the "damaged" macro-volume. Fig.3 indicates that, for a gradual increase in dZ, the choice of a low rate of strain is sufficient .

01

% Jo N X)

-0-1

-02

-0-3 Fig.4. Thermomechanical curves of filled (X) and unfilled (P) polypropylene. -20 C.

Fig.5. As Fig.4, T

_ +40 °C.

222 While investigating the influence of the added filler, it should be borne in mind that the volume of the phase of significant energy consumption decreases with increasing filler content. Fig.3, 4 and 5 indicate that, in dependence on the volume fraction of the filler, as well as on the temperature and rate of strain, the gradual increase in dz may be suppresed or emphasized. Decomposition of the deformation work up to the yield point (W) into parts W^ (where dz is independent of F) and W^ (increasing dZ), reveals a relative increase in

(due to filler) at low temperatures - Fig.6. At elev-

ated temperatures, the addition of filler has the opposite effect. When part of the particulate filler is replaced by fibres, the curve of dz vs. F changes in a characteristic way, especially at the strains just before the yield point. The data in Fig.7 indicate that the higher yield stress of IA (compared with X) is due to processes which produce a small amount of heat. A certain level of deformation seems to be necessary for activation of the reinforcing mechanism of fibres. At higher powder/fibre ratios the cummulation of defects begins at low strain. Restricted mobility of the matrix due to fibres seems to be responsible for this behaviour. Toughening of thermoplastics by dispersed elastomers has a (sometimes undesirable) effect on the compliance increase. Addition of the particulate filler for compliance control requires detailed knowledge composites

of the component interaction. Fig.8 depicts the data for based on a relatively low molecular weight matrix.

£

W, W,

v... 5 mm min-1

100

PX

PX

B

•

P X P X

50

- 20

0

20

AO T(°C)

Fig. 6. The change in W-^/W2 ratio due to addition of filler at different T

o

values,

223 —0-1

0-1

1 o -0-1

.19 1 1 v...5mmmirr • / T o ...-20°C !;

-01

\\

-02

15

5

FtttfN)

AD i ¡AC /A B ' 1 )

-0-2

-0-3

-0-3

Fig.7. Reinforcing effect of glass fibres (IA, IB) added to the PP-CaCC>3 composite.

Fig.8. Thermomechanical curves of PP-EPDM (AB) and PPEPDM-CaCC>3 (AC, AD) composites .

A drop in the yield stress induced by the filler need not be accompanied by any significant increase in the toughness (AC). On increasing the content of the filler, the consequent toughness increase

(AD) is accompanied by a further decrease in the yield

stress. Preliminary experiments using a thermovision apparatus have been carried out, followed by other experiments correlating

"thermo-

elastic" measurements to the acoustic emission signals (5). AE signals of low energy at strains of 10""'"% usually appeared at the minimum value of dz. The tested composites seem to leave the linear elastic strain range at very low tensile stress. Unfortunately, the dynamometers used give no reliable data around the zero force value. The initial drop in dz may be due to an apparent decrease in the material compliance. Further testing is necessary.

224 Conclusion Measuring the changes in the temperature of polypropylene composites during

tensile stressing by amonotonically

increasing load

indicates that (depending on the rate of strain, temperature and composition)

energy dissipating processes may be active in near-

ly the whole range of tensile strain up to the yield point. Even if the level of the energy dissipated by the composite is very low, the observed effects are sufficient evidence for the existence of processes controlling the strength and toughness of the composite macro-volume. More detailed analysis of data is the subject of further research.

References 1. Bucknall, C.B.. 1978. Fracture and Failure of Multiphase Polymers and Polymer Composites. In: Advances in Polymer Science, 27, Failure in Polymers (J.D. Ferry, ed.), Springer, pp.121-146. 2. Chudnovski, A., A. Moet. 1985. J.Mater.Sei. 20, 630. 3. Bever, M.B., D.L. Holt, A.L. Titchener. 1973. The Stored Energy of Cold Work. In: Progress in Materials Science, Vol.17 (B. Chalmers, J.W. Christian, T.B. Massalski, eds.). Pergamon Press. 4. Juska, T., I.R. Harrison. 1982. Polymer Eng.Sei. 22,

766.

5. Hugo, J.. 1982. Thermische u. Schallemissionseffekte bei der zeitabhängigen Deformation von gefüllten Thermoplasten. In: Verstärkte Plaste '82, Kammer der Technik, Berlin.

RIGID STRUCTURAL FOAMS FROM COMPOSITE POLYPROPYLENE/CALCIUM CARBONATE

F. Smejkal Chemopetrol, k.p. Silon Plana nad Luznici, 391 02 Sezimovo Osti II, CSSR

Introduction In recent years rigid high-density foams have been used in new exacting applications, especially in the building industry and agriculture. They are mainly large dimension and thickwall products which serve as construction elements for arrangement in large units, e.g. all-plastic bio-sewage disposal plants, pools, housing units and so on. It is obvious that the demands on quality of structural foams for the above mentioned applications have remarkably increased. High-density foams must have the mechanical properties required and must be suitable from the point of view of quality of their surface appearance. The disadvantage of structural foams, which are prepared by injection molding using chemical blowing agents, is considerable anisotropy of bulk density over the area of the products and thereby variations in their physico-mechanical properties. The quality of high-density structural foams is affected by several variables of which the type and concentration of blowing agents, melt and mold temperature, melt viscosity and injection pressure are most important (1,2). It is mainly the melt viscosity at the decomposition temperature of the chemical blowing agent that shows considerable influence on uniform size and distribution of bubbles in the middle foam layer

(3). If the polymer viscosity is too high

the bubbles will not be able to expand fully; in the opposite case the intercellular walls will rupture and the foam will have an open cell structure. At large-dimension injection molding products with long flow paths different bubble growth rates during the molding filling process will also occur. The most effective method of influencing the formation and growth of bubbles from supersaturated solutions of gases in molten poly-

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

226 mers is nucleation of bubbles. Chemical blowing agents, e.g. azobisformamide, with exothermal decomposition are the most effective. The application of inorganic materials, e.g. silica (4), and finelly-divided metals (5) has already been used, these act as nucleating agents for the bubbles. This has led us to use the composite calcium carbonate reinforced polypropylene for preparing injection molded sheets in the form of high-density structural foams. The aim was to prepare foams with small uniform cells and with the smallest possible anisotropy of mechanical properties over the area of the sheet.

Experimental Rigid structural foams were prepared in the form of sheets, having dimensions 400x400x15 mm and 2000x800x15 mm. The materials used for the study were polypropylenes (Chemickd zavody CSSP, Litvinov) and composites (Chemopetrol, k.p. Silon, Plana nad Luznici). The list and the rheological properties of the polymers used are given in Table 1. As blowing agents Genitron EP-A (Fisons Ind.Chem. UK) (which generates nitrogen, ammonia, carbon monooxide) and sodium hydrogencarbonate (which generates carbon dioxide) were used. All foams were molded in an injection molding machine (CS 7750/ 1200-A, TOS Rakovnik) at a pressure of 12 MPa and at a temperature of 230'C. For the measurement of rheological properties a VP-05 rheometer was used. An Opton machine, Model TGA 10 was used to determine the size of bubbles. The foaming degree of foams is given by the following expression: foaming degree =

where p

p

o " p f x 100 (%)

is the density of polymer, and p_ is the density of foam.

227 Table 1. List and Rheological Properties of Materials Used at 230'C Material

Composition PP PE CaCO, (% wt.)

Taboren PH41C40 Taboren PH41C03 Taboren 2256 Taboren 2257 Mosten 52522 Mosten 52815 Mosten 56935 Taboren 81C40 Taboren PH21C40

60 97

-

1 .235 0. 925 1 .510 1 .100

25

15

30 100

50

20

-

-

-

-

-

-

100 100 60 60

-

Melt flow index _., (21.2 M; (g.cjn g/ (10 min))

40 3 60

-

Density

40 40

2..21 5..43 0..61 0.. 39 4,.21

0. 907 0. 905

14,.28

0. 905 1. 231 1. 240

24..30 8.. 34 0..62

=

T

K.(Y)N

n

log K

0 .. 328 4..062 0.. 384 3..706 0..425 4..025 0..444 4..032 0 .460 .

3.. 551 0..463 3.. 342 0..499 3..107 0..419 3..653 . 0.. 423 4 .011

Results and Discussion Rigid structural foams were prepared as follows: From one mixture of composite or polypropylene (PP) , and a blowing agent about 10 sheets were pressed out ranging from 15 to 50% foaming degree. All the injection molded products were evaluated from the point of view of macrostructure of the middle foam layer and mechanical properties. Considering the scope of the paper, only values for foams with a single foaming degree are presented for illustration in Table 2. In discussion, however, all the measured values were taken into account.

I. Evaluation of the macrostructure of the cellular core Table 2 presents the size of cells in relation to the type of composite and PP for foams with a 25% foaming degree. The dependence of the average size of cells on foaming degree for various foams are given in Fig.1. From the results it is evident, that for the foams made from PP without calcium carbonate, the size of bubbles increases with decreasing melt viscosity (Table 1). In accordance

228

with assumption, the smallest cells are in PP 52-522 and the largest in controlled rheology PP 52815 and 56935, probably about 100 to 200 ym for the same foaming degree. At a small degrees of foaming ranging from 5 to 20% the cells were of ununiform size in all cases. In the range from 20 to 25% foaming the cells achieve the smallest diameters and with further decreasing of density of foams, gradual increase in bubble size is observed. The dependence of the cell size on foaming degree in cellular composites was opposite to the above mentioned cases. The bubbles had the smallest diameters at a low foaming degree and with increasing foaming they slightly enlarged. With the composites a remarkable dependence of the size of bubbles on the melt viscosity has not been observed as in the case of PP. In all cases the defects, i.e. the presence of large, random holes, in the cellular core were found.

Table 2. Average Size of Bubbles for Foams with 25% Foaming Degree

Material

Blowing agent

Size of bubbles x x x min max (ym)

Thickness of skin (ym)

Taboren PH41C40

Genitron EP-A NaHC0 3

404 374

208 167

945 1055

800 1 200

Taboren PH41C03

Genitron EP-A NaHC0 3

201 280

163 240

400 815

1600 171 1

Taboren 2256

Genitron EP-A NaHC0 3

483 353

165 178

348 695

2400 1400

Taboren 2257

Genitron EP-A NaHC0 3

352 350

165 165

827 1040

1400 2400

Mosten 52522

Genitron EP-A NaHC0 3

a 258

a 163

310 613

1 500 1 800

Mosten 52815

Genitron EP-A NaHC0 3

a 440

a 160

590 1517

1 900 2000

Mosten 56935

Genitron EP-A NaHC0 3

219 490

163 1 90

433 1500

21 00 2050

Taboren PH81C40

Genitron EP-A NaHC0 3

430 360

321 210

1158 987

950 1300

Taboren PH21C40

Genitron EP-A NaHC0 3

408 370

392 301

1 680 1510

908 1110

a

Smaller than 150 ym.

229

In the case of the composites with low foaming the most defects were established. The composite PH41C03 showed analogous behaviour to PP without calcium carbonate. The addition of a small amount of calcium carbonate had a remarkable influence on the macrostructure of the foam layer - the bubbles being of the smallest and uniform size of all the cases. The presence of PE in composites Taboren 2256 and 2257 on the quality of foams has not been determined. The shape of the bubbles was also interesting. In PP and the composites with high melt viscosity the bubbles were deformed in the direction of the flow of the melt during the mold filling. At higher foaming degrees the walls of cells were not smooth, but distorted. On the other hand, in PP 52815 and 56935 the bubbles were spherical. In subsurface layers of cellular PP random, large, strongly deformed bubbles were found. These bubbles separate the skin and cellular core. In cellular composites these defects did not occur.

20

30

40

50

f-d-(%) Fig.1. Size of 5281 5 + Taboren NaHCO,;

bubbles versus foaming degree for: (o) Mosten 1% NaHCC>3; (•) Mosten 52522 + 1 % NaHC03; (®) PH41C03 + 1% NaHCC>3; (A) Taboren PH81C40 + 1% (A) Taboren PH41C4 0 + 1% NaHC03.

230

Physico-mechanical properties of rigid foams depend on the macrostructure of the middle foam layer and the proportions of this layer and skin. With the cellular PP, as expected assumption, the thickness of the skin decreases with increasing degree of foaming. It was interesting to find that the foams which were produced by sodium hydrogencarbonate had a thicker skin than foams from Genitron EP-A. Also, foams from controlled rheology PP had thicker skins than e.g. PP 52522. Unlike the values in Table 2, the thickest skin was found in cellular composites. For these composites the values of the thickness of the skin were distorted, because the skin was thought to be the distance from the edge of the cross section to the place where bubbles first occurred. While in the case of cellular PP the skin and the cellular core were sharply restricted, in the cellular composites the bubbles occurred at random in the subsurface layers. The low melt viscosity of controlled rheology PP had a favourable influence on the formation and growth of bubbles in the areas distant from the gate. Consequently, the foams had nearly the same density over the entire

L (mm)

800

Fig.2. Density of foams versus the distance from the gate with the sheets made from: (1) Taboren PH41C40 + 1% NaHCOß; (2) Taboren PH81C40 + 1% NaHC0 3 ; (3) Mosten 52522 + 1% NaHC03; (4) Taboren PH41C03 + 1% NaHC03; (5) Mosten 52815 + 1% NaHC03.

231

area of the injection molded products. While with the cellular PP the density varies at the edges of the sheets, in the case of cellular composites the density increases nearly linearly with the distance from the gate (Fig.2). For the preparation of the foams, blowing agents Genitron EP-A and sodium hydrogencarbonate were used. With the PP the bubbles were remarkably smaller in using Genitron EP-A, mainly for a 20% degree of foaming. In the majority of cases the size of the bubbles could not be measured in view of the measurement range of the microscope Opton. It was surprising that with the cellular composites with sodium hydrogencarbonate the bubbles were smaller than in the case of foams with Genitron EP-A.

II. Physico-mechanical evaluation Physico-mechanical measurement confirmed the well-known observation

f

(g-cm-3)

Fig.3. Impact strength versus density for: (®) Taboren PH41C40 + 1% Genitron EP-A; ( A ) Taboren PH41C40 + 1% NaHC0 3 ; (A) Taboren 2257 + 1 % NaHC03; (•) Mosten 52522 + 1% NaHCC>3; (•) Taboren PH41C03 + 1% NaHC03,- (o) Mosten 52522 + 1% Genitron EP-A.

232 that mechanical properties decrease with the decreasing volume density of the foams. It was surprising that

physico-mechanical

properties of the foams with the same foaming degree which were prepared from the same polymer but with different chemical blowing agents have not been found to be too different

(Fig.3).

The differences in macrostructure of the cellular core due to the type of chemical blowing agent were demonstrated. A question arises in view of the mechanical properties of the foams whether many small uniform bubbles with thin intercellular walls or a smaller number of larger bubbles with thick walls are most effective. The presence of polyethylene at Taboren 2256 and 2257 had a favourable influence on the increasing impact strength. For cellular composites a remarkable drop in modulus of elasticity in relation to the foaming degree was not observed as in the case of the unfilled cellular PP. At the conclusion of this paper it should be emphasized that using composite polypropylene/calcium carbonate for rigid

structural

foams is not suitable for products where uniform properties over the entire area of a part are required. In addition, cellular composites with a max. 25% foaming degree could have been prepared. At higher foaming degrees, however, neither the surface was smooth nor the sheets were well formed. On the other hand, small additions of calcium carbonate have a favourable influence on the macrostructure of the cellular core and thereby on physicomechanical

properties.

References 1. Han, Ch.D., C.A.Villamizar.

1978. Polym.Eng.Sci.

Ij^ 687.

2. Villamizar, C.A., Ch.D.Han.

1978. Polym.Eng.Sci. J_8, 699.

3. Bigg, D.M., J.R.Preston.

1976. Polym.Eng.Sci.

4. Hansen, R.H. W.M.Martin.

1965. Polym.Lett. 3, 325.

5. U.S.Patent

4,317,888.

1_6, 706.

INSTRUMENTED IMPACT STUDIES OF SOME THERMOPLASTIC COMPOSITES

H. Hoffmann, W. Grellmann Institute of Polymeric Materials, Technical University of Leuna-Merseburg, 4200 Merseburg, G.D.R. V. Zilvar Institute of Materials Science, Faculty of Mechanical Engineering, Technical University of Prague, 12135 Prague, Czechoslovakia

Introduction Inorganic particle fillers such as CaCO^ increase the stiffness and heat distortion of some thermoplastics (1). Toughness will decrease unless a proper filler content is chosen. From the data published in refs (2,3) it follows that in case of some composites the Charpy impact strength, a^, reaches a maximum at the characteristic volume content dependent on the particle size and plasticity of the polymer matrix. The aim of this work is to evaluate the impact fracture behaviour of the polyethylene/calcium carbonate (PE/CaCOj) and poly(vinyl chloride). (PVC/CaC03> composites during impact tests using the instrumented Charpy-type impact tester, and to analyze results of the measurements obtained by means of the J-integral, similarly to the case of composites with the fiber fillers (4).

Experiments The matrixes HDPE and PVC, and the filler CaCO^ were commercial products made in G.D.R. The CaCO^ particles had a mean particle size D

equal to 2.5 ym. The energy capacity of the Charpy-type impact - 1

tester was 5 J and its pendulum rate 1.5 ms (4). Dimensions of the specimens exposed to three-point bending were: length L = 80 mm, width W = 10 mm, thickness B = 4 mm, distance between the supports S = 4 0 mm (i.e. S/W = 4) and notch depth a = 2 mm (i.e. a/W = 0.2). The notches had sharp peaks made with a razor blade.

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

234

PE

PE/CaC03 - — DEFLECTION, f

Fig.1. Typical traces of force and displacement: F m - maximum load, fjj, - deflection at maximum load, A c p - area adequate to the energy of crack propagation.

The notch geometry and length were chosen on the basis of a detailed methodic study

(5).

Results and Discussion The Charpy standard impact strength and the J-integral concept The values of the standard Charpy impact strength, a^, and the J-integral have been derived from the load-deflexion plots. The diagrams in Fig.1 are analogous to those for the most typical semi-ductile polymers, described

e.g. in ref.(6). Both elastic

and plastic deformations of specimen took place during the initiation phase, up to the point F . The energy connected with the crack propagation is related to the area A c p which is changing as a function of the filler

content. It is evident that when the

Charpy impact strength, a^, is regarded as the energy stored by the specimen

up to the creation of the crack at the tip of the

notch, the role of the toughness of the composite systems is overestimated when using the standard impact tester, because the propagation of the crack being impeded by the filler. It is therefore more exact to evaluate the impact strength of composites with a ductile matrix in terms of the energy corresponding to the area under the load-deflexion curve up to the point F m

with

235 subseauent propagation of the crack occurring at various load decrease rates and at different displacement characteristics. The values of a^ determined by evaluation of the areas registered by the impact tester with instrumentation reached a maximum, in dependence on matrix type, at a characteristic volume filler fraction, , in accordance with the results published in ref. (2,3). The dependence of a^ versus i(>v in PE/CaCC>3 is shown in Fig.2. The impact strength decreases when

0.2. Its values are

lower by -20% when v = 0.4 at 273 K, and by almost 50% at 300 K than those for PE without filler. Similarly, for the composite PVC/CaCO^ the greatest a^ were measured at

= 0.08. It is dif-

ficult to account for the maxima obtained at characteristic filler fractions on the basis of data supplied by the standard Charpy impact tester. That is evident from Fig. 3, where the effect of on the changes in the maximum force, F , and the corresponding m maximum deflection for both the composites is shown. For PE/CaCO^, (a, ) is determined by the maximum of the dependence of F versus 1 k'max ^ m d> , whereas f decreases continuously with increasing filler conv m centration. At this volume fraction of the filler, the distances between the particles

(approx. 3 ;im according to ref. (7)) are

PE/CaCOj o 273 K

!» I

15

,

•

300 K

***

\

to

1 £

0.5

•

1

l

l

0

0.1

Q2

03

1 OA

VOLUME FILLER FRACTION, f

Fig.2. The dependence of impact strength, a^, on the volume filler fraction, , in PE/CaC0 3 -

236

Fig. 3. Maximum load F m and deflection fjp of PE/CaCC>3 and PVC/CaC03 composites plotted against volume filler fraction . v sufficient, to ensure favorable conditions for the initiation of numerous plastic deformations in microvolumes and, also, for larger plastic macrodeformations of the composite system connected with the formation of a fibrillar structure in the matrix. Therefore, the magnitude of f at T4 =0.2 and at 300 K was still sufficiently m v high. At $ = 0.4 the dispersion of filler in the matrix was not optimal and agglomerates of CaCO^ particles were up to 30 )jm in size (Fig.4). Thus, the f currence of the maximum

values significantly decreased. The oc-

a^ in case of PVC/CaCO^ is not related to

the maximum of the dependence of Fm , but rather to that of f m on . The load increases continuously with increasing filler content up to it - 0 . 2 . The increase in f with increasina filler content v m up to * 0. 1 is due most probably to the formation of fine crazes, which is provoked by a local stress concentration at the boundaries of perfectly dispersed filler particles. We pressume that an agglomeration of CaCO^ particles takes place at 4>v > 0.1, due to high viscosity of the melt. A smaller number of fine crazes is thus created during the impact test. Larger size crazes or defects which appear reguire a higher stress level for their formation and further growth but do not contribute to the development

237

of greater plastic macrodeformation. This hypothesis was examined experimentally by means of electron microscopy. The values of f decrease over (b = 0.1. If the assumptions of linear elastic fracture mechanics (LEFM) were used for processing the results from Fig. 3 and if the critical stress intensity factor, K ^ , were calculated from the F values, the value of K T , would increase up to m la v = 0.17. Therefore, an application of the assumptions based on LEFM could lead to incorrect conclusions. An energetic criterion based on the concept of the J-integral was selected for the evaluation of the fracture processes, inasmuch as a significant plasticity of the specimen was evident from the load-deflection function

(Fig.1). Under suitable experimental conditions, the

results are independent of specimen geometry and dimensions (8). On the basis of a preliminary comparison of various models used for the determination of the J-integral for some polymers (5), the model proposed by Sumpter and Turner

(9) was selected as the

most suitable for the systems studied. It follows from the scheme in Fig.5 that the total deformation energy, A_, Cj is assumed to be divided between elastic energy, U_, hj including uncracked body energy, and plastic component, U p . If 0.2 < a/W < 1, the integral can be defined from the eguation (1)

238

PE/CqC03 Fm

£

-

E

5PVC/CaCO,

LOAD

- >


v = 0.4 remains fragile even at higher temperatures (Fig.6). This is caused by the above mentioned agglomeration of filler particles resulting in a premature lability of the crack due to local stress concentration in the vicinity of the large agglomerates, as well as by limitation of the plastic deformation

o pe * PE /CaC03 (fv =0.2)

173

223 273 323 TEMPERATURE (K)

Fig.6. Critical crack opening displacement 6C(j vs. temperature of PE/CaCC>3 (at the top) and mechanical losses tan 6 at frequency 1 Hz (at the bottom).

240 of the PE matrix in the microspaces between the particles of the filler. High plastic deformations of the matrix are characteristic for low concentrations of CaCO^ and temperatures higher than 275 K, resulting in the formation of fibrills. It is evident from Fig.7 that for the PE/CaCO^ studied no optimum "binding" of the PE matrix to CaCO^ was achieved, in contrast to the results published e.g. in ref.(7). The fibrillation of PE in the boundary layer was observed only in the vicinity of very few of the CaCO^ particles. Most filler particles remain freely deposited in the PE matrix, or the particles are separated from the matrix during the impact test without forming the fibrillar structure in the boundary area. The temperature dependences of the fracture criterion, 2, 11.5 A l ^ , 21.4 CaO, 9.45 MgO, 5.8 FeO, 0.3 Cr_0,, 0.45 MnO, and 0.35 S. d = 0.017 mm and 2 3 _, w 1w = 0.4 mm. Density at 20"C: 2.865 g cm ; 0.95 wt.% APS. Tribological additives were specified earlier (7). Polymerization The polymerization was carried out in a stainless steel vessel under nitrogen, electrically heated and provided with a stirrer, thermocouples and an electronic system to control the adiabatic conditions (6). The diameter and height of the polymer blocks (about 1 kg) were c. 95 and c. 150 mm, respectively. The blocks were left to cool at a rate of about 0.8 K min in the polymerization vessel (6). The products were tested for the yield of the polymer (p), degree of polymerization (P), filler content (w^), tensile and impact properties, dynamic mechanical properties, and morphological structure (6).

285 Results and Discussion The slag (SF) and slag-basalt fibres (SBF) which are much less regular in their form than the glass fibres (GF) affect similarly the adiabatic polymerization of 6-caprolactam and the morphological structure of the polymer. However, unlike the composites with GF, the crystallinity was slightly increased the spherulite diameter followed the same tendency

(Fig.1), while (Tables land 2)

70

I 5

0

I 10

I 15

l 20 25 w. ,wt.7. i

Fig.1. The effect of slag (SF), slag-basalt (SBF) and short glass (GF) fibres on the crystallinity (wa) of poly(6-caprolactam) prepared at T Q = 135"C. Synhydrid/PIC = 0.7/0.7 mol.% for SF and SBF, and 0.3/0.3 mol.% for GF. ( O ) SF, ( Q ) SBF, ( © ) GF with or (C)) without APS treatment.

With SF, contrary to SBF, the rate of polymerization was decreasing

(Table 1) and negatively influenced by a higher Y-amino-

propyltriethoxysilane

(APS) content, which was proved to decrease

the concentration of the growing centres (6). The influence of the different chemical composition of SF and SBF fillers and the polymerization heat losses caused by the heat capacity of fillers (6) also have to be considered. The yield of the polymer

(p) and P value decreased with SF and SBF

content, being affected by APS treatment (6) and a higher activator concentration

(8). However, we can also see a decrease in P in

286

Table 1. Slag and Slag-Basalt Fibres Effect on the Poly(6-caprolactam) Composites3 Fibres content wt. %

P min

wt..%

P seec

Spherulite diameter^ pm

Tensile strength MPa

Shear moduluse GPa

Slag fibres 0 5.7 9.4 15.0 18.1 5.9f

10. 5 13. 2 16. 5 30 .0 41 .0 15. 5

4.7 8.5 15.6 20. 1 8.8f

12. 0 11 .0 12. 0 1 1 0. 23. 5

98.. 5 220 25.0 190 96..7 15.0 96..2 140 7.2 97., 2 85 8.8 115 96..4 8.8 97..5 300 23.5 Slag-Basalt fibres 97.. 5 220 14.7 96..5 180 6.8 96..8 135 3.8 145 96..8 3.6 96.. 6 420 6.0

45.3 36.0 39.3 46.0 -

50.1 51 . 3 48.0 42.0 43.0 54.0

1.18 1 .31 1 .48 1 .64 1 .66 -

1.18 1 .43 1 .44 1 .60 —

(a) Adiabatic polymerization at initial temperature T 0 = 135°C and Synhydrid/PIC = 0.7/0.7 mol.%; (b) polymerization halftime for reaching p = 0.5; (c) degree of polymerization from viscosity measurements in m-cresol solutions; (d) average value; (e) torsional pendulum measurements at 20'C and 1 Hz; (f) Synhydrid/PIC = 0.3/0.3 mol.%.

Table 2. Short Glass Fibres Effect on the Poly(6-caprolactam) Composites3 Fibres content wt. % 0b 5. 1 9.8 14.4 19.3 25.2

V 5

C

P

P

min

wt. %

see c

18..3 16.. 7 18.. 5 19.. 0 19.. 4 21 .5 .

96. 8 98. 0 98. 1 98. 0 97. 9 97. 8

680 480 450 460 410 430

Spherulite diameter0 ym 50. 0 17. 6 16. 6 15. 6 7. 0 4. 5

Tensile strength MPA , without0 with APS 75..9 58., 5 59.. 5 58.. 5 55..8 45..0

75.9 64.8 59.3 58.0 51.0 56.5

(a) Adiabatic polymerization at T 0 = 135°C and Synhydrid/PIC = 0.3/0.3 mol.%; (b) average value from 10 polymerizations; (c) cf. Table 1; no APS treatment; (d) APS 0.2 wt.%.

287

composites with GF both with

(Table 2] or without APS treatment

(6). Similarly to the previous results (6), the effect of SF and SBF fillers on the molecular mobility of the matrix estimated using dynamic mechanical measurements was characterized by a linear increase in the shear modulus

(Table 1). More pronounced (a)

relaxation process and a shift of the main

(B)

relaxation process

towards lower temperatures evidenced an increase in the concentration of the unreacted monomer. Table 3. Poly(6-caprolactam) Composites with Tribological Fillers 9

Filler

(wt.%)

None d f

Friction -r e coef.

Wear*3 . -1 pm km

p wt.%a

P

96.8

770®

0..36

33

Notched Tensile impact strength 0 . c strength1kJ m~2 MPa 5.4

76.5

97.6

230

0..18

200

3.5-4.8

G

(10) f

97.5

400

0 .14 ,

307

3.0-4.0

50

G

(25) g

96.7

425

0.. 4

230

4.8

29

-

0..17

45

-

-

-

0 .09 .

16

-

-

gel

0..18

4

4.0-6.0

40

435

0. 1 5

9

5.0

26

MOS 2

(10)

G/MoS 2 /PTFE/Gb (5/5/10/10) g G/B/O

(15/5/5} g

G/MoS,/0

(15/5/5)

(Type SP-27M) G/MoS 2 /0

(18/9/9)

(Type SP-29M) h

9 6 9

45-50

(a) Adiabatic polymerization; (b) rotatory motion tests against steel, no lubrication (4); (c) dry test specimens; (d) TLA/PIC = 0.3/0.3 mol.%, T 0 = 135°C; (e) Synhydrid/PIC = 0.3/0.3 mol.%, T 0 = 135'C, P = 680; (f) Synhydrid/PIC = 0.6/0.6 mol.%, T C = 135"C; (g) Synhydrid/HDI = 0.65/0.33 mol.%, T 0 = 145'C; (h) Synhydrid/PIC = 1.3/0.9 mol.%, TQ = 145"C; G graphite; 0 engine oil SAE 90; Gb glass beads 0.15 mm; B bronze powder below 0.2 mm.

The polymer yield in the composites prepared with tribological additives

(Table 3) corresponds to the values typical of the

polymer without fillers. However, the P values were lower primarily due to a higher initiator and activator content

(8) and a negative

288

SBF •GF~ "SF

1

0-05

0

0-1

v. f

F i g . 2 . E f f e c t of the slag ( S F , 0 ) , slag b a s a l t ( S B F , © ) a n d g l a s s (GF,t)) f i b r e s o n t h e s h e a r m o d u l u s (G) o f p o l y ( 6 - c a p r o l a c t a m ) a t 2 0 ° C a n d 1 H z f r e q u e n c y . F u l l l i n e s p r e d i c t e d by t h e o r i e s d e s c r i b e d e a r l i e r (6); v f in t h e f i l l e r v o l u m e f r a c t i o n (cf. T a b l e 1).

a c t i o n of

impurities

(4), p o l y m e r s was

lower

smaller

a broad distribution

of

toughness impact

in d i a m e t e r .

in the

(a)

30 t o 4 0 ° C ;

shear modulus

changed

of

the

resistance

the composites

and

e v a l u a t e d by

The

morphological

the

some

process tensile

decreased

fillers

in a good a c c o r d w i t h or the H a l p i n - T s a i

M and SP-29 M

the p r e d i c t i o n

theory

(10)

account to

strength,

The

6). T h e

of

(Fig.2).

notched

increased.

The

(cf. T a b l e

3)

stiffness

a method described earlier

found to be

with

molecular

temperature

(cf.

(1.15 G P a ) .

and

low

strain-at-break

SP-27

with

crystallinity

irregular

and oil used w i t h

Nielsen

(9)

polymer.

main relaxation

little,

The

As expected,

the u n f i l l e d polymer value

the composites

In a g r e e m e n t

occurred.

consequently,

and temperature

strength

equalled

surfaces.

(2 t o 4 vim), o f t e n

in the c o m p o s i t e s

for a decrease a value

filler

in m - c r e s o l

in c o m p a r i s o n w i t h u n f i l l e d

formations were products

on the

insoluble

the

(6)

of

was

Kerner-

289 References 1. Miller, R.E. 1963. U.S. Patent 3,344,107. 2. Hedrick, R.M. and P.A. Tierney. 1964. U.S. Patent 3,418,268. 3. Brit.Celanese Ltd. 1964. French Patent 1,360,518. 4. Strouf, O., B. Cásensky and V. Kubánek. 1985. Sodium Dihydrido-bis(2-methoxyethoxo)-aluminate (SDMA). Elsevier, Amsterdam, p. 207. 5. Bukac, Z. and J. Sebenda. 1974. Czechoslovak Patent 169,258. 6. Horsky, J. and J. Kolarik. 1985. Acta Polymerica 36^, 220 and 225. 7. Horsky, J., L. Holy and F. Bása. 1980. Czechoslovak Patent 220, 449. 8. Wichterle, 0-, J. Sebenda and J. Králícek. 1961. Fortschr. Hochpolym.-Forsch. 2, 578. 9. Nielsen, L.E. 1974. Mechanical Properties of Polymers and Composites. M. Dekker, New York, p. 388. 10. Halpin, J.C. and J.L. Kardos. 1976. Polymer Eng. Sei. 16, 344.

ANIONIC POLY(6-CAPROLACTAM) COMPOSITES POLYMERIZED IN ROTATING MOULDS

J. Horsky Central Research Institute, Skoda Works Plzen, 316 00 Plzen, Czechoslovakia

Introduction The activated anionic polymerization of 6-caprolactam in rotatory mould is a processing method widely used in the manufacture of tubes (1), pulley discs and wheels (2), calender rollers (3) and other products. The polymerizing system consists either of one lactam component or of a combination of two lactams which can form more layers of the polymer with different qualities (1). The rotary processing of poly(6-caprolactam) can be highly useful in the production of bearing bushes made of polymeric composites with either solid or liquid lubricants. These systems have been extensively studied in connection with a new initiator system based on the organic complexes of aluminium salts (4). Suitable processing methods with stationary moulds (5) were developed and advantageous sliding properties of the material were determined (6). A very serious technical problem of a centrifugal lactam polymerization in situ with high density fillers is their sedimentation in early stages of polymerization. As the lactam melt viscosity is too low to prevent the sedimentation of the fillers, it is necessary to control the beginning of the tube moulding process. Equation (1) derived in ref.(7) describes, under some simplified conditions, the behaviour of a spherical particle (with a radius R and density p^) in a surrounding medium (characterized by the viscosity n and density p ) situated in the centrifugal mould (a speed of rotation n per min, the outer and inner turning radii r 2 and r^, the rotation time t in s): ri = 8 I T 2 P.2(pf- p m ) t n 2 /9 X 3600> and of the block 6. Through an opening 6-caprolactam and into its melt other constituents as Synhydrid, Desmodur H and fillers were added. At the initial temperature T Q = = 135°C, the polymerization was started and proceeded on heating 8 under adiabatic conditions

(3) for about 2 min. After an increase

in viscosity indicated by the stirrer power input, while the temperature reached c. 137°C

(about 5 to 10% conversion of the mono-

mer to polymer), the polymerizing mixture was transported into the rotary mould 9 by pressure of gaseous nitrogen. The mould

(inner

diameter 112 mm, length 485 mm) was electrically preheated J_0 (through an outer cylinder jacket) to 125°C and during the time of filling it was at a standstill. When the polymerizing mixture was fully transferred into the mould

(within 1 min), the rotation

started with 800 r.p.m. for about 10 s, and then the heating was set up to 180°C at a heating rate of 5 deg min~1, while the speed

293 0 W

A

Fig. 1. Polymerization and rotating mould appliances for composite tubes processing of rotation was readjusted to 500 r.p.m. By means of thermocouples in the cylinder jacket V\_ and at the inner wall Y2, the heating was controlled during the polymerization (8 to 10 min) and annealing (30 min) periods, which were followed by a spontaneous cooling at a rate of 1.3 K min down to 100°C; then the heating was switched off, while the mould was still rotating at 500 r.p.m. At about 100°C the motor J_3 was stopped and the mould was removed from the frame bearings 1_4- Consequently, the tube was taken off the mould after the barrel heads 15 were dismantled. The produced tubes were tested for the yield of the polymer, degree of polymerization, moisture and filler content, and size of the morphological formations (8).

Results and Discussion The study resulted in a development of several composite materials, of which the type of SP-265 (Table 1) was proved to be the best fitting the working conditions of selflubricating liners of cushion cylinders for undercarriages of electric locomotives. The adiabatic prepolymerization as well as the polymerization in the rotary mould proceeded with an appropriate rate and were characterized by a high yield of the polymer (97 to 98%). However, the yield was lower (88 to 93%) in the end parts of some tubes close to the edge of the mould, where small heat losses were inevitable.

294 Table 1.

Characteristics of the Cylindrical Tubes with Outer and Inner Diameters of Length 108/70/470 mm Respectively, Polymerized in the Rotating Mould

Polymerization number SP-234,la Ca Cb Oa SP-249,la Ca Cb Oa SP-265,la Ca Cb Oa

Polymer yield wt. % 97.2 97. 8 97.6 97.3 96.9 96.9 96.9 93.1 88.0 96.6 96.5 95. 6

Moisture content wt. % 2. 7 2. 5 2. 5 2.7 2. 0 2. 0 2. 0 4. 6 10. 5 2. 2 2. 2 2. 6

P

Filler content, wt.% Introduced

0.1545 -

2094 1275 -

360 294 -

417

20.16 20. 16 20. 16 20.16 20.00 20.00 20.00 20.00 25.04 25.04 25.04 25.04

Measured 18.4 27.2 17.5 19.2 22.1 24.0 20.0 24.3 22.5 54.4 15.8 28.8

SP-234, SP-249: Synhydrid/Desmodur H =0.65/0.65 mol.%, filler: graphite CR-5; SP-265 : Synhydrid/Desmodur H = 1.0/1.0 mol.%, filler: graphite/MoS2/oil= 15/5/5 wt.%; I, C, 0 specimen extracted either from the inlet, central, or outlet part of the tube; a, b outer and inner surface of the tube, respectively. The average degree of polymerization, P, was very high in some places of the composites prepared with graphite. Usually, the block polymer with 25 wt.% graphite content produced by the adiabatic polymerization of 6-caprolactam initiated by 0.65 mol.% of Synhydrid and activated by 0.33 mol.% of Desmodur H at T q = 144"C was characterized by P = 425 or with the Synhydrid and Desmodur H concentrations 0.7 mol.% and T q = 135'C, a polymer without any filler resulted with P = c. 200. The increased P indicated in Table 1 was most likely caused by a partial loss of the activator in chemical reactions with unspecified pollutants that are always present, as highly inert system conditions at the pilot plant were difficult to be granted. The morphological formations examined in a microscope under polarized light appeared irregular in the crystallites shape and distribution with composites of the graphite type (Fig. 2a), whereas 'the structure of composites with a combination of fillers (Fig. 2b) was much more uniform. The moisture content after a 12-month storage in the air at 50 to 80% R.H. and 20 to 25"C was comparatively low and uniform in the composite mass, though the values were higher in some end parts of

295

Fig. 2. Morphological structures: (a) SP-234 composite tube with 20 wt.% graphite content, (b) SP-265 composite tube with graphite/ MoSp/oil fillers (cf. Table 1). the tubes, probably due to a higher error of measurement caused by the increased monomer content (Table 1). A comparison between the samples taken of the outer and inner surfaces of the products revealed a certain sedimentation of the filler in the central part of the tube length due to the hydrodynamic conditions in the filling period of the mould or during the first steps of polymerization in it, because the prepolymerized mixture was retained in this place a little longer. According to an estimate of the sedimentation time calculated by means of Eq.(1) and using approximately known viscosity at the start of rotation and early polymerization stages (Table 2), a small sedimentation can be expected for the graphite particles within a distance of 1 to 1.5 mm from the tube surface, and also a more severe sedimentation of MOS 2 particles in the whole tube wall. Because the real mixture conditions (e.g. n or p ) were surely different from those assumed, the measured differences in the filler content were not so pronounced (Table 1), and were acceptable for the operational use of the tubes, due to the machining allowance. The tubes indicated a shrinkage 2.9 to 3.2% in the axis and 3.0 to 4.5 in the diameter direction.

296 Table 2.

Sedimentation Time for Graphite and M0S2 Filler Particles in a Rotating Mould

n

Delay in Approx. Supposed Sedimentation time for sevthe mould monomer to fluid vis- eral ^ / r ^ (in mm) , s min ^ polymer con- cosity s version (10-12) m Pa s 56/54.5 56/54 56/55 Graphite CR-5 (a) 10 800 5 26 13 19 9 180 500 15 76 115 154 21 360 500 229 346 50 63 463 Molybdenum disulfide (b) 5 800 10 9 0.14 0.1 1 80 500 15 21 0.6 0.8 360 500 50 63 1 .8 2.5 (a) R = 2.5 x 10~6 m , p f = 2. 27 x 1 O 3 kg m~ 3 , P m = 0.97 * 1 0 3 kg m - 3 . (b) R = 2 x 10~5 m, pf = 4.8 x 1 0 3 k g m - 3 , p m = 0. 97 * 103 k g m _ 3 . p f and p m values from refs. (13,14). -

, %

References 1.

Geisler, N. 1979. British Patent 2,045,151 A.

2.

Fräser, M. 1981. German Patent DE 3,143,401 A-1.

3.

Küsters, E. 1964. German Patent 1,214,865.

4.

Strouf, 0., B. Cäsensky, V. Kubänek. 1 985. Sodium Dihydrido-bis (2-methoxyethoxo)-aluminate (SDMA). Elsevier, Amsterdam.

5.

Horsky, J., L. Holy. 1980. Czechoslovak Patent 218,678.

6.

Horsky, J., L. Holy, F. Bäsa. 1980. Czechoslovak Patent 220,449.

7.

Tanford, Ch. 1961. Physical Chemistry of Macromolecules. J.Wiley & Sons, New York. pp. 258, 324, 364.

8.

Horsky, J., J. Kolafik. 1985. Acta Polym. 36,

9.

Horsky, J., L. Holtf. 1984. Czech.Patent Applic. No PV 10,029-84.

10. 11. 12. 13. 14.

220 and 225.

Vieweg, R., A. Müller. 1966. Kunstoff Handbuch, Vol.6, Polyamide. Carl Hanser Verlag, München, p. 677. Simünkova, E., J. Zelinger, V. Kubänek, J. Krälicek. 1973. J.Appl.Polym.Sei. V7, 1387. Gabbert, J.D., A.Y. Garner, R.M. Hedrick. 1983. Polym.Composites 4, 196. Dean, J.A. 1973. Lange's Handbook of Chemistry. Mc Graw Hill, New York. pp. 4-35, 82. Kohan, M.I. 1973. Nylon Plastics. J.Wiley & Sons, New York, p. 465.

MECHANICAL PROPERTIES OF SOFT PVC-TEXTILE COMPOSITES

J. Pig^owski, M. Koz^owski Institute of Organic and Polymer Technology, Technical University of Wroclaw, 50-370 Wroclaw, Poland

Introduction For obvious security reasons, the conveyor belts used in coal mines should be slow-burning. In this respect, there is a trend throughout the world to replace rubber belts by those made of soft poly (vinyl chloride)

(PVC). Application of specially selected types of

plasticizers further decreases fire hazards. The best for this purpose are aryl or alkylaryl phosphonate plasticizers. Also essential is the low viscosity of the PVC-plasticizer plastisols. This parameter decides whether or not a type of plasticizer can be used. Low-viscosity plastisols are suitable as they enable quick impregnation of usually thick textile carriers. A plasticizer meeting such requirements has recently been developed in the Institute of Organic and Polymer Technology of Wroclaw Technical University. It is a phosphonate plasticizer modified with epichlorhydrine based on phenol and hexylphenol. The mechanical properties of PVC modified with this plasticizer and the rheological characteristics of plastisols have been described by Dul et al. (1). The composites studied in this work consist of polyester

(PET) and poly-

amide (PA) unwoven cloths of similar basis weight (ca. 150-200 g/m2) impregnated with the above mentioned plasticizer.

Experimental Pastes: 5 2 g of the phosphonate plasticizer, 80 g of emulsion PVC (Meronyl, France) and 1.6 g of organotin stabilizer

(Ergoterm BTGO,

Poland) were mixed together in a high-speed mixer. Cloath coating: Unwoven cloths were coated from one side with the paste and the system was pregelled at 378 K for 10 min. The other

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

298

side of the cloth was then coated. The both side coated cloth was gelled at 438 K for 15 min. The resulting composites had 1.4 mm in thickness. Mechanical properties of the cloth PVC composites and the adhesion of the soft PVC to the cloths and test films were measured using an Instrom tensile strength machine. For the adhesion measurements, strips having 1 cm in width were used, while for the strength measurements the normalized dumbbells were prepared.

Results and Discussion The tensile strength measurements carried out at the rate of deformation 10 mm/min led to results strongly dependent on the direction of deformation with respect to the orientation of fibers. The highest tensile strength was observed when the deformation was parallel to the direction of either weft or warp. The tensile strengths measured for these two mutually perpendicular directions

Table 1. Mechanical Properties of Composites Carrier

Force N

Strength MP a

385 90

44.4

32

60 ° i

70 57 337

10.4 8.1 6.6 38.9

115 153 100 32

i*

290

33. 4

27

290 75 92

33.2

33

30 ' 45"

8.6 10.7 10.4 33.2 30.5

110 97 100 43 45

Direction il

PET

PA

30° 45 °

60° 1 1*

90 290 265

Elongation a "5

*Samples pretreated with an adhesion improving preparation

299

were almost the same for polyester cloth and exactly the same for polyamide cloth. Identical were also the relative elongations at break. In the directions ranging from 30° to 60° with respect to the warp direction, the tensile strengths were found to be several times lower than those measured for the parallel direction, while elongations at break were substantially higher. In order to examine the effect of adhesion upon the mechanical properties of the composites, the adhesion of plasticized PVC to polyester and polyamide was studied for the systems soft PVC-polymer film. The results are shown in Table 2. The breaking force (calculated per unit length) increases with the rate of deformation. It is worth noticing that the breaking force is equivalent to the specific work of adhesion and, for the quasi-equilibrium conditions, it can be identified with the thermodynamic work of adhesion. The relationship F = f(V) can be expressed in the form of power law F = kV n

(1)

where F is the breaking force per unit length, V is the extension rate, and k and n are constants. It is believed (2) that k is related to the size of defects existing in the material, while n describes the relaxation behaviour of the system. Since the properties of the carrier (PET, PA) are very

Table 2. The Results of Stripping of Soft PVC from Test Films Carrier

Extension rate V mm/min 10

PET

50 100

500

PA

Force F

F = kV

N/cm 0. 35 0.84 1 .00 1 . 40

10

0.08

50

0.16

100

0.22

500

0.28

F = 0.18V

0. 35

0.32 F = 0.0 4V

300 different from those of soft PVC, one should expect the relaxation properties of the PVC component to play a major role. In this case, the values of n for both carrier polymers should be the same, as it was observed indeed. The differences ink can probably be ascribed to the differences in the interfacial tensions between the two pairs of components. The interfacial tension can affect the

*

magnitude of the defects in a joint. Much higher forces needed for PET system than for PA system (cf.Table 2) seem to confirm this conclusion, at least indirectly. Attempts have been made to improve adhesion of the soft PVC to the polymer by treating it with vinylidene chloride-acrylonitrile-itaconic acid terpolymer (3), resorcinol and surface active agents. Applying a deformation of 10 itun/min, the values of F = 0.66 and 1.7 N/cm were obtained for PET and PA systems, respectively. These data are several times higher than those presented in Table 2 for the non-pretreated samples. The mechanical properties of the pretreated cloth composites, however, decreased in similar way. To explain this observation adequately, some adhesion measurements with several composites have also been made. Like in the case of films, a higher adhesion of soft PVC to polyester unwoven cloth was observed than that to polyamide cloth. The absolute values were much higher than those for films (Table 3), since mechanical adhesion (trapping of the coating on the developed surface of cloths)

Soft

F

\f\NN, distance

Fig. 1. Scheme of the delamination process.

301

Table 3. The Results of Stripping of Soft PVC from Unwoven Cloth Characteristics

1

F, N/cm

11.0 1 .5 2.0 1.7

1 Q r mm mm ^ - A l , mm

— 1

2 8.5

8.0

1 .5 1 .7 1.4

2.7 3.5 3.0

2 7.7 2.7 3.2 2.7

1 Untreated samples; 2 samples pretreated with adhesion improving preparation. For definition of 1, 1 0 and ^ see Fig.1 and Eq.(2)

prevailed. In an extreme case, because of the presence of voids in the cloth, through which the PVC layers on the opposite sides came into contact to form a cohesive joint, the adhesion measured approached the value of cohesion. As it follows from Table 3, the adhesion decreases after a preparation layer has been applied. This explains the worsening of mechanical properties of the pretreated cloth composites. Microscopic examinations provide an answer to the question why the preparation enhances the adhesion to solid smooth material and fails for cloths. It has been found that the preparation layer applied to the cloth by dipping left behind a thin film preventing the PVC plastisol to penetrate into the cloth, thus eliminating the mechanical adhesion. During the delamination experiments, a characteristic periodical change in force F was observed (Fig.1). The increase in force occurred when the material between warp fibers was subjected to deformation. Then, the adhesion-cohesion interaction system had to be overcome. A drop in force required for delaminating was observed for the PVC layer lying over warp fibres (just only adhesion to be overcome). From the cycle length, 1, it was possible to calculate the dimensions of repeating structural elements in the cloth, 1

o

=

(v

r / v p' ( 1 / 2 )

(2)

where v r is the rate of deformation and v ing paper. For rigid joints

is the rate of record-

= 1 q . In the case of the composites

studied, however, the PVC layer was stretched to a considerable extent, while the one-side coated cloth did not change dimensions;

302

consequently ^

> 1 Q was observed. The value of

can be correc-

ted for elongation under stress. By taking the value of elasticity modulus at 100% elongation (9.9 MPa (4)) and the mean stress from Table 3, it was estimated that II o - A1 was close to real 1o value. Concluding, one may state that the orthotropic composites studied in this work are characterized by a strong dependence of their mechanical properties on the direction of force applied. It has been found that a pretreatment of cloths aiming at the improving adhesion reduces the penetration of plastisols inside the cloths, thus making the mechanical properties of the composites worse. The preparation improving adhesion of the soft PVC either to polyester or to polyamide might be applied to fibres prior to forming unwoven cloths.

References 1. Dul, M. , M. Koz^owski, J. Pigjiowski, VJ. Rospond. 1 985. PrzemysZ Chemiczny _64, 135. 2. Gul, W.E. 1971. In: Structure and Resistance of Polymers. Chemistry, Moscow, pp. 189-213. 3. Pat.PRL 242051 (1983). 4. Koz^owski, M., J. PigXowski. 1984. Report No.21, Properties of PVC - Phosphonate Plasticizer Systems. Technical University of Wroclaw.

HYDROPHILIC AND THERMOFORMABLE SILICONE RUBBER COMPOSITE

P. Vondräcek, J. Hrudka Department of Polymers, Institute of Chemical Technology 166 28 Prague 6 J. Sulc, P. Lopour Institute of Macromolecular Chemistry, Czechoslovak Academy of Sciences, 162 06 Prague 6

Introduction Silicone rubber is one of the few suitable elastomers for biomedical use. It has fairly good mechanical properties sufficient for most biomedical applications, such as implants, tubing or artificial organs, and it is relatively inert and stable when implanted in the body. A typical characteristic of silicone rubber is its hydrophobicity

(1).

On the other hand, hydrogels, typically lightly crosslinked polymers of 2-hydroxyethyl methacrylate

(HEMA), are hydrophilic ma-

terials capable of rather massive water uptake. They exhibit about 70% swelling by volume in water and an outstanding tolerance in the living tissue (2). They are attractive as biomaterials because they are similar to the body's own highly hydrated composition. It is possible to cover silicone rubber or other hydrophobic elastomers with a biocompatible hydrogel coating by various techniques (3) and thus to combine good mechanical and processing properties of rubber with the surface hydrophilicity. Recently, experiments have been performed to prepare interpenetrating networks of a hydrogel and an elastomer. These materials, based on either silicone (4) or polyether urethane (5) rubber, absorbed water like a hydrogel, but had mechanical and processing properties superior to the latter. We report here the preparation and properties of a composite material with similar characteristics based on conventional peroxide-crosslinked silicone rubber filled with a particulate hydrogel filler: further details are to be described elsewhere.

Polymer Composites © 1986 Walter d e Gruyter & Co., Berlin • New York - Printed in Germany

304 Materials and Methods A commercial peroxide-cured silicone elastomer (Lukopren G 1000, Lucebni zavody Kolin, Czechoslovakia), basically dimethyl siloxane rubber, having a viscosity-average molar mass of about 500 000 and containing a small percentage of vinylmethylsiloxane units was used as a rubber matrix of the composite. Lightly crosslinked poly(2-hydroxyethyl methacrylate) particles of a BET specific surface area ranging from 4 to 10 m 2 /g, containing 3 to 10% of methanol extractable sol fraction were prepared by slurry polymerization under intensive mixing in toluene of the HEMA monomer containing 1 to 3% of ethylene dimethacrylate as a crosslinking agent. The preparation of the poly(HEMA) particulate filler was based on a patented polymerization procedure (6) followed by mechanical disintegration of dried poly(HEMA) agglomerates in a mortar and sifting the product through a sieve with the mesh size of 0.04 mm. Mixes of silicone rubber filled with 20, 50 or 100 phr (parts per one hundred parts of rubber), i.e. 16.7, 33.3 or 50 wt.% of poly (HEMA) particles, were prepared by a conventional rubber mixing technique on a two-roll mill at room temperature. Bis(2,4-dichlorobenzoyl) peroxide in the form of 50% paste in silicone oil (1.2 phr) was added as a curing agent for silicone rubber. Vulcanized sheets of 1 mm thickness were prepared from the mixed compound by compression moulding in a laboratory press at 110°C for 15 minutes. If not specified otherwise, all data published in this paper were obtained with poly(HEMA) particulate filler of the BET specific surface area 4.4 m 2 /g and methanol extract 8.9%, containing 1.5% of ethylene dimethacrylate. Swelling behaviour of the composite was studied in various solvents which do not swell silicone rubber but do swell the hydrophilic poly(HEMA) particulate filler, such as water, ethanol and ethylene glycol monoethyl ether ("Cellosolve"). Swelling studies are carried out on 1 mm thick round samples of an approximately 0.2 g weight. Samples are swollen in duplicates at 20"C for two months, change in weight being recorded. The maximum swell recorded during this period is used as a measure of swelling in this study. Stress-strain properties in tension were measured on standard dumbbell test pieces, 1 mm thick, at various temperatures (20 - 150"C) at cross-head speed 1.67 itun.s-^ (Instron tensile tester). The

305

stress-strain properties were also studied with samples swollen to the maximal swell achieved in the swelling period of two months in the solvent used. Stress values were then related to the swollen cross section area. To test the thermal shape memory of the composite, strips cut from the sheets were stretched to a 100% extension at temperatures ranging from 70 to 150°C and then cooled in the stretched state. A residual elongation (set) was determined. The predeformed samples were then swollen in water and the residual elongation was measured in one-hour intervals for the total 5 to 8 hours of swelling..

Results and Discussion Typical swelling data of the composite are given in Table 1. The degree of swelling was expressed as a change in volume by swelling, i.e.(V-Vo>/Vo,where V is the volume after swelling and V Q is the initial volume. Data in Table 1 show that the composite swells in water, i.e. the hydrogel particulate filler renders silicone rubber hydrophilic. Table 1 . Composite Total (A) and Poly(HEMA) Filler (B) Change in Volume (%) by Swelling in Various Solvents at 20°C as a Function of Poly(HEMA) Content Poly(HEMA)

H20

content, phr

A

0 20 50 1 00

0.9 16 20 32

EtOH

B

A

_

1 27 74 73

Cellosolve

B

1.1 34 68 102

A 257 246 236

1 .3 51 105 1 60

B

_

387 380 370

The degree of swelling of the silicone rubber-hydrogel composite in water and polar organic solvents increases with hydrogel filler content as expected

while the swelling degree of poly(HEMA) par-

ticles in the silicone rubber matrix rather moderately decreases with their concentration in the composite. This is probably a result of mechanical constraint to swelling of a hydrogel particle by other swollen particles in its vicinity at higher hydrogel concentrations .

306

Swelling of poly(HEMA) particulate filler in the composite seems to be higher than that of the hydrogel itself. For example, poly (HEMA) increases its volume by 68% when equilibrium swollen in water (2), while the change in volume of the composite related to its poly(HEMA) content shows higher values (see Table 1, B values), especially at lower hydrogel concentrations. Similarly, the reported (2) volume changes by swelling in ethanol and Cellosolve of poly(HEMA), i.e. 122 and 274%, respectively, are lower than our results of poly(HEMA) swelling in the silicone rubber matrix. It is not clear what is the reason for this apparent increase in the particulate poly(HEMA) swelling in the composite. It might be caused by some vacuole formation and solvent entrapping at the interphase during swelling. The swelling study showed a slight reduction in the total composite swell with an increase in the BET specific surface area of the hydrogel particles. For example, the volume change by swelling in water dropped from 32 to 22% when the BET area of the hydrogel filler increased from 4.4 to 9.5 m 2 /g in the composite containing 100 phr of such filler. Investigation of the mechanical properties of the composite showed that the poly(HEMA) filler enhances the silicone rubber modulus, expressed as stress at 100% elongation, and tensile strength in the dry state at 20"C in dependence on the filler concentration as shown in Table 2, where the stress-strain data in tension for the composite filled with various concentrations of poly(HEMA) filler are summarized. This effect is more pronounced with poly(HEMA) Table 2. Stress-Strain Properties of the Dry (A) and Water Swollen (B) Composite at 20'C Poly(HEMA) content phr 0 20 50 100

Stress at 100% elongation, MPa A B

Tensile strength ,MPa A B

Elongation at break, % A B

0.27 0.45 1 .01 ND

0. 30 0. 76 1 . 26 2.20

130 200 250 ND

ND 0.21 0.17 0.09

ND 0. 38 0. 42 0.56

ND 250 350 420

ND - not determined particles of higher values of specific surface area, which is usual

307 for any other particulate filler. Tensile strength and 100% modulus values of the composite are reduced by the hydrogel filler swelling close to those of an unfilled silicone rubber vulcanízate The stress at 100% elongation drops even below the original value of the silicone rubber matrix. The higher the hydrogen filler content, the greater the modulus and tensile strength reduction by swelling. On the other hand, elongation at break is regularly increased by swelling of poly(HEMA) particles within the rubber matrix. Heating the composite to elevated temperatures ranging from 70 to 150"C resulted in an effect on mechanical properties very similar to that of poly(HEMA) swelling as shown in Table 3, where data are summarized for the composite filled with 50 phr of the hydrogel particles. Table 3. Effect of Elevated Temperature on Mechanical Properties of the Composite Filled with 50 phr of Hydrogel Particles Temperature, 'C

20

70

90

110

150

Stress at 100% elongation, MPa 1.01 0.53 0.50 0.38 0.20 Tensile strength, MPa 1.26 1.24 1.17 1.12 0.49 Elongation at break, % 250 450 480 600 630 Modulus and tensile strength values of the composite decrease with temperature to a value close to that of the initial silicone rubber matrix (see Table 2). At 150°C (T of poly(HEMA) is about 100"C) the 100% modulus drops below the value of matrix similarly as in the case of swelling. Elongation at break increases with temperature. The effects brought about by swelling or heating the composite may be attributed to a decrease in the polymeric filler modulus. The reinforcing effect of the filler is reduced as the filler modulus decreases close to or below the modulus value of the rubbery matrix. Figure 1 allows us to compare all the described effects, namely the reinforcing action of poly(HEMA) particles and influence of heating and solvent swelling on the: tensile behaviour of the silicone rubber-hydrogel composite. Summarily, dry poly(HEMA) particulate filler has a reinforcing effect in silicone rubber at temperatures below its T g , while it increases elongation at break at temperatures above T , or in the swollen state.

308

ELONGATION

Fig.1.

(7.)

Stress-strain curves in tension for silicone rubber filled with 50 phr poly(HEMA) particles at various conditions, compared with tensile behaviour of unfilled silicone rubber. 1 unfilled silicone rubber vulcanizate; 2 composite in a dry state at 20°C; 3 composite in a dry state at 150°C; 4 water swollen composite at 20°C.

Non^swollen poly(HEMA) softens when heated above its Tg. The softening of the hydrogel particulate filler also renders the silicone rubber composite thermoformable. When the material is deformed, e.g., stretched at a temperature higher than 100°C and then cooled in the deformed state, a rather high residual deformation results, depending on the hydrogel filler concentration, as shown in Table 4. This residual deformation caused by changing the shape of filler particles at higher temperature and its following fixation by cooling is completely reversible by swelling in water or in polar organic solvents or by heating the composite material above T^ of poly(HEMA). The elimination of the residual elongation is illustrated in Table 4. Preliminary experiments showed that the addition of pyrogenic silica to the composite reduces swelling only slightly and allows to prepare composite with improved mechanical properties. This is illustrated in Table 5 where the properties of silica reinforced (20 phr Aerosil 130, Degussa) silicone rubber composite are compared with those of the non-reinforced one. Both materials contained 33 wt.% of poly(HEMA) particulate filler.

309

Table 4. Residua-l Elongation after 100% Elongation at 1 35 "C and its Elimination During Swelling in Water at 20 "C Time of swelling, h 20 phr poly(HEMA) 50 phr poly(HEMA)

0

1

2

3

4

5

45 85

30 45

20 25

10 15

5 7

0 3

6 -

0

Table 5- Effect of Silica Reinforcement on Properties of Silicone Rubber-Poly(HEMA) Composite Containing 33 wt.% of Hydrogel Particles Property Tensile strength at 20°C, MPa Tensile strength at 150°C, MPa Volume swelling in water, %

Non-reinforced Reinforced (20 phr Aerosil 130) 1 . 26 0.49 20

2. 60 1 .06 17

Conclusions A composite material consisting of crosslinked silicone rubber matrix and lightly crosslinked poly(2-hydroxyethy1 methacrylate) particles was prepared by a conventional rubber processing technique. The composite swells in water and polar organic solvents. The new type of composite material thus combines hydrophilicity and thermoformability of the hydrogel filler with the processibility of the basic elastomeric material. It can be deformed at elevated temperatures (above T of poly(HEMA)), and keeps the obtained shape if cooled down in the deformed state. The composite shaped in this way resumes its original form if it is again heated or allowed to swell in a polar solvent, i.e. the thermally preshaped composite possesses a shape memory. Since both initial materials are biocompatible, the authors suppose that the new composite material may be suitable for biomedical applications, especially for medical devices which could be fixed in the body by swelling, or which could be inserted into a body cavity in the thermally preshaped form, where it would swell by body fluids to the original shape. A patent is pending (7).

310

References 1. Braley, A.S. 1971. Elastomers for implantation in the human body. Rubber Chem. Technol. _44, 363-380. 2. Wichterle O. 1971. Hydrogels. In: Encyclopedia of Polymer Science and Technology, Vol. 15 (H.F.Mark, N.G.Gaylord and N.M.Bikales, eds.). J.Wiley and Sons, New York. pp. 273-291. 3. Hoffman, A.S. 1975. Hydrogels - a broad class of biomaterials. In: Polymers in Medicine and Surgery (E.L. Kronenthal, Z. Oser and E. Martin, eds.). Plenum Publishing Corp., New York. pp. 33-44. 4. Vale B.H. and R.T. Greer. 1982. Ex vivo shunt testing of hydrogel - silicone rubber composite material. J. Biomed. Mater. Res. ¿6, 471-500. 5. Dror, M., M.Z. Elsabee and G.C. Berry. 1979. Interpenetrating polymer networks for biological applications. Biomat., Med. Dev., Art. Org. T_, 31-39. 6. Chromecek R. and I. Gavrilovä. 1979. Czechoslovak Patent 138 856. 7. Sulc J., P. Vondräcek and P. Lopour. 1985. Czechoslovak Patent Application PV 3955-85.

EFFECT OF POLYMER MATRIX ON THE EFFICIENCY OF MICROCAPSULATED FLAME RETARDANTS

T.V. Popova, R.P. Stankevitch, M.S. Vilesova, N.A. Khalturinskii, Al.Al. Berlin Institute of Synthetical Polymeric Materials, USSR Academy of Sciences, Moscow, USSR

Introduction Polymers and polymer-based composite materials have found wide application in various branches of national economy. However, their use is limited due to high flammability. One of the ways to reduce flammability of polymeric materials is the introduction of flame retardants into polymer matrix. However, the introduction of additives into polymer compositions may deteriorate physico-mechanical properties of the material due to poor compatibility of the additives with the polymer. Moreover, during the processing of the material liquid additives may be evaporated and extracted from polymer matrix by solvents. So, it is reasonable to use certain flame retardants in microcapsules, that is, included in a polymer shell. Using an appropriate material for a microcapsule shell, one can stabilize and even improve physicomechanical properties of a polymer composition since in this case a flame retardant also plays the role of a modifier. The application of microcapsulated flame retardants (MFR) allows us to obtain materials with a working surface enriched with a flame retardant. This increases the flame resistance of material without considerable changing its physicomechanical properties. The use of MFR simplifies their introduction into a polymer solving the problems of volatility of low-boiling liquids, chemical activity and thermodynamical incompatibility of liquid flame retardants with a polymer matrix. The processing conditions (temperature, pressure) of polymer compositions, containing microcapsulated additives, should be such as to keep the shell of microcapsules unbroken during the processing. The shell material as regards the polymer composition components

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

312 should be either inert or interactive with the polymer itself or with the composition components. This strengthens the microcapsule shell and makes stronger its connection with the polymer. The mechanism of MFR action has its own features, which also depend on the polymer matrix they are incorporated into. To clarify the mechanism of MFR and the effect of a polymer matrix on the efficiency of their action, we studied the influence of a flame retardant type shell material and a microcapsule size on the temperature of the opening of a capsule itself and of that included in a polymer matrix. We also compared the efficiency of flame retardants incorporated into polymer matrix as plasticizers and in microcapsules.

Experimental Results and Discussion Trichlorethylphosphate, tri-(2,3-dibrompropylphosphate), carbon tetrachloride, freon-114B2 and freon-318B2 were used as additives. Polyvinyl alcohol, gelatin and polyurea served as a shell for microcapsules. In Table 1 are given the temperatures of MFR "opening"

(T

).

From the Table it is clear that at the heating rate of 10 K s T

coincides with the temperature at the beginning of intensive

shell (sh) material decomposition

'T170 ° C

>200 ° C

17.0

39.0

Without MFR

Foam Plastic

GED+PEPA 335"C

ED-20+PEPA 340"C

16.0

22.0

Freon-114B 2 T, = 4 7 "C b Shell-gelatin

-

-

25.0

32.0

-

-

37.5

43.5

60.0

60.0

T, . ,,=70'C d ( sh) Freon-318B 2 T, = 98 " C b Shell-PVA T

d(sh)=230'C

CC1 4 T, = 77'C b Shell-PVA

18.0

40.0

22.0

45.0

T , , ,-=230°C d ( sh) Tri(2,3-dibrompropyl)phosphate

-

15 mass %

22.0

5%

25.0 15%

Trichlorethylphosphate

23.0

46.5

15 mass %

-

23.0

5?

28.1 15?

their inhibiting action. As well as non-microcapsulated phosphates high-boiling phosphorous-containing MFR slightly improve flame resistance

(oxygen index increases by 5-7%) due to larger char

formation on the surface of the burning polymer. In the case of reactoplastics with higher temperature of decomposition as compared with T, . , . low-boiling MFR in the highly

320

overheated state "explode" long before the beginning of intensive polymer matrix decomposition. One can observe strong effect of dispersion. The size of dispersed particles and the rate of their carrying away depend on the force of the "microexplosion", which is larger the larger the difference between T^gjj) a n d of a flame retardant. Moreover, the effect of dispersion of reactoplastics and, thus, the efficiency of inhibiting action depends on the type of low-boiling flame retardants, included in a shell. So, the inhibiting action of carbon tetrachloride, which is an inert diluter of flammable degradation products of polymer matrix in the combustion zone, is weaker when it is incorporated into polymer as a plasticizer than that of bromine-containing freons, which are gas-phase inhibitors in the flame and increase char formation. However, during the application of these flame retardants in microcapsules one can observe the opposite dependence of the efficiency of their action. The most effective among low-boiling MFR under consideration is carbon tetrachloride. At its 5% content in an epoxide composition the values of oxygen index reach 60% whereas the introduction of 10-12% allows us to obtain the material non-flammable in the atmosphere of pure oxygen due to strong dispersion. Fig.4 shows the influence of a microcapsule size of different lowboiling microcapsulated liquids on the efficiency of the inhibition of the epoxide composition combustion at the MFR concentration of 7 mass %. From the Figure it is clear that the size of microcapsules considerably influences the values of the oxygen index of compositions for all low-boiling liquids considered. All of them have extreme values of oxygen indices, corresponding to the optimal size of microcapsules (each microcapsulated low-boiling liquid has its own optimal size of microcapsules and its own value of the oxygen index maximal for the given concentration). Optimal size and concentration of low-boiling MFR depend on thermochemical and thermophysical characteristics of the capsulated liquid as well as on thermomechanical characteristics of the polymer matrix and the shell material. So, MFR inhibit the combustion of polymer material more effectively when dispersion of the burning material takes place, i.e. when the boiling temperature of a liquid flame retardant is much

321

01, %

200

d, |jm

300

Fig.4. The influence of a microcapsule size, d, and the type of MFR on the effective inhibition of the combustion of the epoxide composition ED-20 cured by PEPA: 1 - CC1 4 ; freon-114B 2 - MFR content - "7%. freon-318B 2 ? 3 lower than T ^ g ^ )

(flame retardant being overheated) and T 3( ma i-)

is higher than T, . , a (sh) place.

at which the process of dispersion takes

In conclusion it should be mentioned that the application of MFR is an effective and prospective way of reducing the flammability of polymer composite materials.

PROPERTIES OF PROTEIN MODIFICATIONS COVALENTLY LINKED TO PARTICLES

0 . Zemek, L. Kuniak Institute of Chemistry, Centrum of Chemical Research, Slovak Academy of Sciences, Dûbravskâ cesta 9, 842 38 Bratislava 1. Novék, D . Berek Institute of Polymers, Slovak Academy of Sciences, Dûbravskâ cesta 11, 842 36 Bratislava, Czechoslovakia

Introduction Biopolymers can be immobilized either by their crosslinking with bifunctional agents or through a covalent linkage to the particles of a carrier acting in the system as a composite. This can be mediated by a spacer of different length and type (1). We report here the effect of the particle total surface and the length of spacers on the amount of a protein linked. Properties of the composed systems in respect to the protein modifications prepared in the crosslinking are compared.

Experimental Materials. The silicagels were prepared by polymerization of acidified solutions of silicates of alkaline metals followed by washing in water and organic solvents and heated to 750 °C according to (2). The silicagel preparations were characterized according to (3). ¿*-Aminopropyltriethoxysilane {IT-APT), thiophosgene, 1,6-diaminohexane and 1,12-diaminododecane were Ferak (Berlin West) products; 1,2-diaminoethane (Oenapharm, GDR); 1,8-diaminooctane (Merck, Darmstadt, FRG); 2,4,6-trinitrobenzen14 sulfonic acid was from Serva (Heidelberg, FRG), (U— C)L-valine (0,22 MBq.mol" 1 ) and (U— 1 4 C)L-cysteine;(0.22 M B q . m o l - 1 ) were the products of the Radioactivity Centre (Amersham, England). 131 I human serum albumin (HSA, the starting specific radioacti-

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

324

vity 0.5 M B q . m g - 1 ) and

131

I albumin egg (0.7 M B q . m g - 1 ) was dona-

ted by the Institute of Nuclear Research, Radioisotope Production and Distribution Centre, Otwock, Poland. The diisothiocyanates used in the experiment were prepared from the corresponding diamines and thiophosgene according to (4). Methods. ,^-APT-silicagels. The silicagels (2 g) were silanized according to (5) with If-APT (25 % concentration in toluene (50 ml)). The reaction product was suction-filtered using a sintered-glass filter, washed with toluene, ethanol and air-dried Polyisothiocyanates of silicagels: The H 2 N-derivatives of silicagels (2 g) were reacted with cC. co -diisothiocyanates in chloroform (40 ml) at 90 °C for 8 h . Determination of - N H 2 groups. Procedure according to (5). df-APT-silicagels (20 mg) were added into a pH 8 borate buffer (0.2 mol.l ) containing 2,4,6-trinitrobenzene sulfonic acid mm

A

fl

(5 mmol .1" ) and reacted at 40 C for 1 h. The undissolved portion was centrifuged and the supernatant (0.1 ml) added to vali_A ne (20 mmol.l ) in 1 % trichloroacetic acid (0.1 ml). After a 2 h reaction at 40 °C, 0.5 m o l . l - 1 HC1 (5 ml) was added. The -NHg content was calculated from calibration graph for valine at 410 nm. Reaction of silicagel isothiocyanates. The silicagel isothiocyanates (10 mg) were suspended in 0.1 mol.l borate buffer (pH 14 8.2; 1 ml) and stirred with (UC)L-valine, or (U- 1 4 C)L-cyste131 ine or I albumin human or egg (50, 100, 1000, 3000 yug) at 35 °C for 5,10,15 and 20 h. The derivatives of silicagel were then washed with the above mentioned borate buffer (5 ml), the respective L-valine, L-cysteine (5 ml, 10 mg.1 ) or albumin —1 —1 (5 ml, 1.5 mg.1 ), 20 mmol.l NaCl, borate buffer, and finally it was dried in a vacuum drier. The incorporated radioactivity was measured with a Packard counter using a toluene scintillation liquid SLX-31. Protein crosslinking. Human serum albumin (1 g) or albumin white egg (1 g; Serva, Heidelberg, FRG both of them) were dissolved in a borate buffer (0.05 mol.l ; pH 9, 10 ml) and crosslinked with 1,2-diisothiocyanatoethane (C 2 ), 1,6-diisothiocyanatohexane (C g ) or 1,12-diisothiocyanatododecane (C 1 2 ) i n the ratio from 1 : 1 up to 1 : 50 (C/P m Table 1). The reaction proceeded at 30 °C for 5 h (6). After washing the unreacted

325 compounds, the proteins (1 g) were characterized and compared with those linked to silicagels. Calculation of the spacer lengths. The maximum distance between the Si atom of -APT and the reaction center of the linked spacer already made, i.e. up to carbon atoms of -NCS was calculated from geometric parameters (bond length, valence angles) in the respective spacers (7).

Results and Discussion Properties of the crosslinked proteins are summarized in Table 1. As we have not succeeded in any of the experiments in finding the residual crosslinking agent in washings after the reaction, we suppose that all the amount of the diisothiocyanates reacted with the proteins. As we have not demonstrated any affinity of 14 the crosslinked proteins to (U— C)L-cysteine, we suppose that all diisothiocyanates reacted in the crosslinking reaction and not in the side reactions. Table 1 . Characterization of the Crosslinked Human Serum Albumin (a) and Egg Albumin (b)

Cross 1 inking agent

C/P

1.0

0.2 0.02

1.0 0.2 0.02

1.0

0.2 0.02

V, ml .g

1

b

a 5.8 6.4 15 .2

3.1 3.5

6.5 6.9

3.2 3.7

17.1 7.1 7.8 18.2

10 .2 4.1 4.9 12 .3

8.2

C/P: the weight ratio of diisothiocyanates and protein; V : bed volume

326

As demonstrated in Table 1 the length of a crosslinking agent used does not influence substantially the swelling ability of the crosslinked gel formed. The C/P ratio in the reaction mixture is of substantial importance. The lowering of diisothiocyanates concentration is followed by a marked increase in the swelling bed volumes in water. The silicagels prepared according to (2) are characterized in Table 2 . Table 2 . Characterization of Silicagel Amins Derivatives of silicagel

Total n^.g- 1

Silicagel 1 -R Q Silicagel 2 -R o Silicagel 3 -R Q Silicagel^-R Q

-NH2

surface ymol.g

30

72 .3

156 370 652

173.5 261.2 375.7

R q =-NH(CH 2 ) 3 -Si(0H 2 )-0-for Table 2 and 3 According to results summarized in Table 2 the lowering of the total surface of the particles is followed by an unproportional decrease in the - N H 2 content. The - N H 2 content is, however, proportional to the extent of the reaction with diisothiocyanates. This reaction was monitored using (U— 1 4 C)L-valine and (U— 1 4 C)L-cysteine (Table 3 ) . The amount of the linked albumin increases with the increasing total surface of the silicagel particles. For the interaction with both albumins spacers prepared from 1,6-diisothiocyanatohexane,1.82 nm were unequivocally favoured. Therefore the highest amount of the albumins linked was in the case of silicagel^ and the Cg spacer. Amount of the human serum albumin linked to the silicagel in this case reached 373 mg , g _ 1 . Our finding is in agreement with refs. (1,8). Capacity of the silicagels for HSA binding is higher than that for A egg binding probably due to accesibility of pure HSA protein chains in the reaction. (A egg is a glycoprotein). When we used the diisothiocyanates of various length either as agents for protein crosslinking or as spacer for the protein binding to the silicagel

327

particles, their effect was more pronounced on the phase boundary of the silicagel particles than in the crosslinking reaction. Table 3 . Characterization of Silicagel Xsothiocyanates Derivatives of silicagel

L-valine L-cysteine -1 |i(nol.g

HSA A egg -1 mg .g

v -1 ml .g

Spacer nm

1 .2 1 .38 1 .22

1 .34 1 .82 2 .63 1 .34 1 .82 2 .63

67 .5 68 .1 66 .9

12 .6 32 .8 16 .5

11 .8 19 .9 17 .7

85 .7 87 .5 83 .1

144 .2 148 .2

72 .1 156 .4 98 .0

43 .1 95 .7 68 .5

1 .41 1 .52

150 .2 156 .1 153 .5

231 .1 238 .5 293 .3

143 .0

1 .55

253 .1 178 .1

67 .8 95 .3 73 .2

1 .58

1 .34 1 .82

1 .56

2 .63

Silicagel 4 - R 0-°2 245 .1 Silicagel^ - R O " C 6 247 .4 Silicagel^ - R O " C 1 2 235 .2

356 .2 358 .5 365 .7

285 .1 373 .4 309 .0

157 .1 195 .2 167 .2

1 .62 1 .7 1 .6

1 .34 1 .82 2 .63

Silicagel^ - R 0 ~ C 2 Silicagelj R C - O" 6 Silicagel^ R C - O- 12 Silicagel 2 R C - O" 2 Silicagel 2 - R O " C 6

24 .6 24 .7

Silicagel^ - R O " C 2

Silicagel 2 - R O " C 1 2 Silicagel 3 R C - O~ 6 Silicagel^ R C - O" 12

24 .4

145 .3

1 .48

Contrary to the effect of the crosslinking agents, the silicagel particles as composites in the system substantially suppress the swelling bed volume of the total composed system as well as that calculated from the protein amount linked. This is of great importance when reaction zones of reactor based on immobilized enzymes operating in aerosols are to be constructed.

References 1. Flemming C . A . Gabert, H . Wand, 3 . Zemek. 1983. Collection Czechoslov. Chem. Commun. 48, 184. 2 . Zemek a., I. Novdik, D . Berek, L . Kuniak. 1984. Czechoslov. Pat. 211 176

328 3 . Brunauer S., P.H. Emmett, E . Teller. 1938. 0 . A m . Chera. Soc. 60, 309. 4 . Wand H . 1978. Acta Biol. Med. Germ. 37, 501. 5. Wand H., M . Rudel., H . Oautzenberg. 1978. Z . Chem. 18, 224. 6 . Zemek 0., L . Kunlak, T.L. Yourkshtovich. 1985. Makromol. Chem. Suppl. 9, 2 2 7 . 7 . Pople O.A., O.L. Beveridge. 1970. Approximate Molecular Orbital Theory. McGraw-Hill, New York, p. 110. 8 . Cuatrecasas P. 1971. 0 . Agr. Food Chem. 19, 600.

THE REVERSIBILITY OF HYGROTHERMAL EFFECTS IN FIBRE-RESIN COMPOSITES

G.Pritchard, S.D.Speake School of Industrial, Organic and Polymer Chemistry Kingston Polytechnic,Penrhyn Road, Kingston-upon-Thames, Surrey, KTl 2EE, England

Introduction Absorption of water by fibre-resin composites results in significant changes in their mechanical properties. In practice, however, reinforced plastics products are rarely used in circumstances requiring continuous immersion in water, and periodic drying often occurs. It is therefore of interest to know the extent of reversibility of property changes on drying, and the factors affecting material recovery. Recently, Kasturiarachchi and Pritchard (1) have reported the extent of recovery of certain properties of unidirectional glass reinforced epoxy laminates. Scanning electron microscope evidence was presented to show extensive interfacial bond restoration on careful and prolonged drying. There was also complete recovery of shear strength. Signs of some permanent alterations in the matrix, notably cavitation, were evident. This paper is concerned with reversibility in glass reinforced polyester laminates.

Materials An isophthalic polyester resin, Cellobond A283/270, was supplied by B.P. International.

It was used in the cast, reinforced form

and also as the matrix for glass laminates.

Polymer Composites © 1986 Walter d e Gruyter & Co., Berlin • New York - Printed in Germany

330 Hand lay up was carried out by Thermoset Laminates Ltd. of Leigh, Lanes., U.K., using (a) unidirectional fibre bundles, held together with transverse stringers, and (b) crossed plies arranged to give - 45° configuration. The laminates were cured at room temperature for 24 hours and postcured at 80°C for 5 hours. They were then cut into strips 200 x 25 x 3.25 - 0.2 mm thick, with fibre orientations 0°, 10°, - 45° and 90° to the strip axis. Unidirectional fibre content was 48% w /w (29% v/v) and the - 45° strips had glass content 40% w /w (2 4% v /v).

Mechanical Properties Tensile dumbells were machined from cast sheet to the geometry of B.S. 2782 Part 3 (1976) Method 320 B, using a computer controlled milling machine. Laminate strips were used without further shaping. Mechanical tests were performed with two conventional screw-driven tensile machines (Instron 1114 and Nene M3000).

Conditioning, immersion and drying Samples were first dried over silica gel at 4 5°C for two weeks and then over phosphorus pentoxide for a further week. They were then weighed, totally immersed in distilled water in baths controlled to - 0.5°C , and re-weighed periodically. They were cooled in cold water prior to mechanical testing. Redrying was carried out by storing for 14 days at 45°C over silica gel and 14 further days at 45°C over phosphorus pentoxide. This period is not quite sufficient to reach constant weight, but results in a very close approximation. The temperature of 45°C was chosen to circumvent microcracking.

331

Subsequent examination Specimens were examined by optical microscopy to detect disc cracks and to assess fibre-resin adhesion. Laminates were examined for light transmission in a visible light spectrophotometer. The change in absorbance at 600 nm of each laminate was assessed by measuring absorbance in ten different locations chosen at random, and taking the average value. There is no change in the absorbance of unreinforced resin samples at 600 nm on immersion in water.

Results and Discussion Degradation processes in cast resins Re-drying unsaturated polyester resins after prolonged immersion in hot water results in a final weight less than that of the original, because of the leaching of extra-network material. The water absorbed during immersion not only swells the network, but replaces much of the non-bound organic matter. The nature of the leached substances has been discussed in a previous publication (2). The theory of Flory (3) predicts that there will inevitably be some oligomeric species in the reaction products of difunctional condensation reactions, and assuming random copolymerization of fumaric and phthalic acid units, some of these oligomers will be uncrosslinkable. In addition, the sol fraction will include unreacted crosslinking agent, catalyst diluent, and catalyst residues. It follows that the re-dried resin cannot be the same substance as a virgin sample. Consequently, the true weight of water absorbed at a given time t, (M^) is given by the sum of the observed weight increase M ^ an 2 samples in Figs, in and 11.

As indicated above, these samples are

only poorly cured, resulting in numerous free chain ends: it is to the motion of these chain ends that we attribute this peak.

Their absence in the low n

series II samples suggest segments longer than n=l are required.

It is noted

359 the 6" peak is not seen in the mechanical spectra, for which it would be expected at ahout 0°C: for either the series I samples or the phenoxy resin. In the latter case the concentration may be too low: in both types of samples it may not be mechanically active although the postulated mechanism would suggest it should he. Results similar to the above were also seen for the -y, B and B 1 peaks for the Series I curing conditions using Epon 828 (Fig. 13). increases

In particular 6'

with cure at 150°C, with 6 and 8' increasing during postcure.

In

addition to the expected peaks, a B" peak is seen which was attributed above to free chain ends.

Although Epon 82ft has some monomer ( n a v e = 0.1) and linear

polymerization may be the initial reaction during curing, the 6' peak increases

with cure rather than decreases as in Fig. 1?.

At this time we do

not know if its presence is related to some unexpected feature of the curing process, or if the molecular assignment is incorrect.

Further clarification of

the origin of this peak would be desirable.

Moisture fects

absorption

(exposure to atmospheric humidity)

primarily

af-

the 3 peak in Epon 828 (Fig. 14), both increasing its strength and

reducing its temperature; the effect is attributed to a plasterization effect in the matrix.

Considerably smaller effects are seen on the 3' and B" peaks.

6' stays at essentially the same temperature, increasing slightly in strength, whereas B" decreases in temperature and increases in strength.

The rise at

near room temperature is due to ionic impurities and varies greatly from sample to sample.

• w e b iSk~,2ocrc u. •acre kr Ik •«rei»

•

WATER EXPOSED

°

DRY

? o K

9

z

.. ^•••MMI»* ito

TEMP fx) 220

265"

Figure 13. Effect of curing conditions on catalyst cured Fpon 828 TSD spectra. Series I.

105

'

fis

'

TEMP n o 2 2 0

260

Figure 14. Effect of moisture absorption on TSD spectra of catalyst cured Epon 828. Sample cured at 1R0°C for 15 h, followed by 4 h at 200°C.

360 Anhydride cured samples.

Only mechanical

the anhydride cured Epon

relaxation spectra were obtained for

samples as a function of degree of mixing and

ratio of mono- to difunctional

curing agent.

In the 1.00% difunctional

anhydride cured sample two low temperature mechanical

relaxation peaks are seen

in addition to the possible presence of the y peak at -130°C (Fig. 15, sample S2).

The peak at -80°C is again attributed to the cross-links

new peak is seen at - ? 0 ° C . agent segments

(B) whereas a

This peak is attributed to motion of the curing

(Fig. 1): it decreases in height with increasing

"curing agent" concentration.

mono-functional

It is noted the B peak remains nearly

constant

in temperature and height with varying ratios of curing agents, the mobile junction unit is apparently about the same regardless of whether it is .ioined to another chain or terminates in a mono-functional

anhydride unit.

not of primary concern here, there is a large decrease in Tg with monofunctional mixing

agent concentration, as would he expected.

Although

increasing

Increased degree of

(Fig. 16) results in an increase in sharpness of both peaks and a slight

increase in the S peak position, presumably due to the more

homogeneous

environment for the mohile units involved in the two relaxations.

Accompanying

the increase in homogeneity of cross-linked density is significant

improvement

in room temperature ultimate properties

(Table 3).

Summary

In summary, we have observed, and attributed to the motion of specific molecular units, 5 relaxations in catalyst cured epox.y resins at temperatures below room temperature.

They are tabulated in Table 1.

An additional

relaxation was seen in the anhydride cured sample and related to motion of the anhydride units incorporated into the network.

Of concern now is correlation

of this relaxation behavior to room temperature impact properties of both the resulting network polymer and as incorporated in composite

systems.

Acknowledgement

This research was supported by the National

Science Foundation Polymer

Mitsubishi

Flectric Co. (T.T.) and Chung Shen Institute of Science and

Technology

(C.S.W.)

Program,

361

-160

-120 Tamp. *C-40

0 '

40

Figure 15. Effect of degree of mixing on tan 6 of anhydride cured Epon R?fl, standard cure. S - l . RO rpm: S-2, 300 rpm: S-3, 2000 rpm: a l l for 5 minutes.

0 40 160 -120 Tamp, • C " 4 0 Figure 16. Effect of di/monofunctional anhydride r a t i o on tan « of Fpon 82R. S-2. 1/0: S-5, 8/1: S-6, 4/1: S-7, 2/1 mole r a t i o . Standard cure.

References 1.

Takahama. T., P.H. G e i l . 1.9R2. J. Polym. S e i . . Polym. Phys. Ed.. 20. 1979.

2.

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3.

Takahama, T., P.H. G e i l . 19R2. J. Pol.ym. S e i . . Polym. Lett. Ed.. 20. 453.

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Takahama T. 19R2. M.S. Thesis, U n i v e r s i t y of

6.

Wu, C . S . . 19R5. Ph.D. Thesis. University of I l l i n o i s , papers in preparation.

7.

Chen A.. 1.9R5. M.S. T h e s i s , U n i v e r s i t y of I l l i n o i s , paper in preparation.

R.

Pangrle, S . . 19R5. M.S. T h e s i s , University of I l l i n o i s , paper in preparation.

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S i n n o t t . K.M.. 1962. SPE Trans.,

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10.

Andrews, R . n . . T.J. Hammack. 1965. J. Polym. S e i . R, _3_, 655.

11.

Matsuoka, S . , Y. I s h i d a . 1966. J. Polym.Sei. C, _14. 247.

12.

Williams, J . G . . 1979. J . Appl. Polym. S e i . . 23_, 3433.

13.

Van Hoorn, A.. 196R. J. Appl. P o l y m . S e i . , _12, R71.

14.

Pogany G.A.. 1970. Polymer, _11_, 6 6 .

15.

Oammont. F . R . , T.K. Kwei. 1967. J . Polymer S e i . A-2, _5, 761.

16.

Delatycky, 0 . . J.C. Show, G. Williams. 1969. J. Polym. S e i . A-2. 7, 753.

362 17.

Arridge. R.G.C., J.H. Speak. 197?. Pol.ymer,

443.

1«.

Ochi, M., M. Shimbo. 1976. Nipon Kagaku Kaishi.

19.

Ochi. M., T. Takahama, M. Shimbo. 1979. Nippon Kagaku Kaishi, _5, 662.

20.

Mijovic. J.. L. Tsay. 1981. Polymer, 22, 903.

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Prath, E.H.. M. Robinson. 1963. J. Polym. Sci.. Part C, _3, 65.

22.

Erath, E.H., R.A. Spurr. 1959. J. Polym. Sci. 3S. 391.

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24.

Cuthrel 1, R.F.. 1967. J. Appi. Polym. Sci.

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Kenyon A.S., L.E. Nielsen. 1969. J. Macromol. Sci.-Chem. _3,

26.

Kreibich, U.T., R. Schmid. 1975. J. Polym. Sci., Symp. 53, 177.

27.

French. n . M . , R.A.H. Strecker, A.S. Tompa. 1970. J. Appi. Polym. Sci. 14. 599.

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Strecker. R.A.H.. 1969. J. Appi. Polym. Sci.

29.

Lewis, A.F.. 1963. SPE Trans. 3, 201.

30.

Lewis, A.F., W.R. Ramsey. 1966. Adhes. Age 9^, 20.

31.

Ousek, K., J. Plestil , F. Lednicky. 1978. Polymer

32.

Matyi R.J.. D.R. Uhlmann. 1960. J. Polym. Sci.. Polym. Phys. Ed. _18, 1053.

33.

Racich, J.L., J.A. Koutsky. 1976. J. Appi. Polym. Sci. 20. 2111.

34.

Mijovic. J.S.. J.A. Koutsky. 1979. Polymer j!0. 1095.

35.

Mijovic. J.S., J.A. Koutsky. 1979. J. Appi. Polym. Sci. _23, 1037.

36.

Morgan, R.J., J.E. O'Neal. 1977. J. Mater. Sci.

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Woohnsiedler. H.P.. 1963. J. Polym. Sci. Part C, 3, 77.

38.

Mijovic, J., L. Tsay. 1981. Polymer 22^, 902 .

39.

Aspbury. P.J., W.C. Wake. 1979. Br. Polym. J. 1 U

40.

IJhlmann, D.R., J.B. Van der Sande. 1979. In Discuss. Faraday Soc. 68, p. 129.

41.

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44.

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>

1495.

161 . 43.

697

?439.

393 .

1966 .

17.

~ 3503.

THERMALLY STIMULATED DEPOLARIZATION OF RADIATION CURED UNSATURATED POLYESTER RESIN-GLASS MICRONODULES COMPOSITES

Z. Jelcic, F.

Ranogajec

"Ruder Boskovic" Institute, Zagreb 41000, pob 1016, Yugoslavia

Introduction The modification of unsaturated polyester

(UP) resins by rein-

forcing additives has attracted considerable interest in recent years, as a mean of improving the mechanical properties. A largely empirical approach was typical of initial work on these additives. One of the most popular additives is glass. Glass particles applied as reinforcing fillers of UP resin matrices greatly improve their mechanical, electrical and thermal properties. Cured UP resin is suitable insulating material. Its conductivity monotonically varies during curing, generally decreasing with time and tending to a constant value when curing is complete. Thermally stimulated depolarization

(TSD) currents of the thermo-electrets

formed on cured polyester systems modified with different amounts of glass micronodules have been measured to study the effect of polymer-filler interface on the carrier-transport mechanism. Interfaces and surfaces are always involved with composites and frequently, if not invariably, exert a controlling influence on their electrical properties. Dielectric polarization and depolarization studies were performed on radiation cured samples. Two techniques were used. First, the poling

(polarization) current

was measured at fixed times, starting from the application of the electric fields (1.5 kV/cm), as a function of temperature. Using the other technique

(TSD), the partially cured sample was polar-

ized at a temperature which was higher than the transition temperature, cooled down with the field on, and reheated at constant rate with the field off, and the shortcircuit current coming from the sample was recorded as a function of temperature.

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

364 Experimental The used polyester resin system industry

(H 201) supplied by Chemical

(HINS), Novi Sad, Yugoslavia, H 201 is based on maleic

anhydride, propylene glycol and orthophtalic acid, and contains 40 wt.% of styrene. The commercial filler was polycellular

(P25)

glass micronodules filled with a gas from SOVITEC, Glaverbel, Belgium. The materials were thoroughly mixed and then cured, both at room temperature, by gamma irradiation

(dose rate 3.6

kGy/h)

in the coaxial aluminium cells for TSD measurements. Due to very low density of Microcel P25 and low viscosity of the resin

(600

centipoises), samples of the radiation cured UP resin composites were found to undergo phase separation into two layers. The top layer appeared to consist essentially of a very high volume fraction of glass beads. The bottom layer was more homogeneously

cured

and insoluble in benzene.

Results and Discussion Figure 1 shows the TSD spectrum of nonfilled UP resin cured by several gamma radiation doses. It is characterized by

several

partly overlapping peaks which will be designated by y, 8, a,a' , and p in order of increasing temperature. The TSD spectra for filler UP resins of varying crosslinking density and composition

(cured by 4 kGy)

(from 6.25 to 25 vol.% of glass beads) are plotted

in Figures 2 and 3. It can be seen that the TSD spectrum of control sample

(cured UP resin) reaches lower values than the cor-

responding composite TSD spectra. Poling current experiments (Fig.4) give a peak characterized by the same position, and relative height as the low temperature TSD peaks

(Fig.5). Figure 5

summarizes TSD transition temperatures of UP resin composites containing glass beads. These data closely follow the corresponding neat resin transition temperatures shown also in Figure 4. Small differences observed between the transition

temperatures

probably are not due to changes in the chain mobility or in the free volume. This could be probably caused by heterogeneous

struc-

ture of the systems. Also, the transition temperatures have been determined from the slight humps on the TSD curve. The character

365

F i g . 1 . T S D s p e c t r u m of r a d i a t i o n c u r i n g of H 2 0 1 .

F i g . 2 . T S D s p e c t r u m of r a d i a t i o n c u r i n g of H201

+ P25

(25 v o l . % ) .

366

F i g . 3 . T S D s p e c t r u m of H201

+ H 2 5 of v a r i o u s g l a s s c o n t e n t

F i g . 4 . P o l i n g c u r r e n t of H201 (4 kGy) in v o l . % .

+ P 2 5 of v a r i o u s g l a s s

(4 k G y ) .

content

367 T/K 300

250

200

kGy

Fig.5. Transition temperatures by TSD.

of the peaks (Fig.5) is close to that predicted for the intrinsic relaxations: their locations, nearby independent of the composition, are close to that of nonfilled UP resin and their amplitudes are functions of the filler concentration. This probably involves the charge carriers previously accumulated at the phase boundaries have been detrapped by the molecular motion occurring during the glass transition. No very well pronounced peak of individual transitions has been observed for cured composites, except of that for the & transitions, but slight humps at TSD curves, corresponding to the dipolar peaks, have been observed. Note that as the filler content increases the TSD and poling current are enhanced. This increase in conductivity with increasing filler content is due to states forming in the gap. It can also be seen that the TSD maximum appears a few degrees lower than the poling current maximum. This probably results from the fact, that, due to the existence of a distribution of relaxation times, the polarization increase occurring during cooling concerns mainly the slow relaxation processes. The TSD spectra of filled and nonfilled

368 UP resin were analysed and several distinct transitions were observed which were identified as the two glass transitions a and B, two liquid-like transitions a' and p, and a segmental transition y.

Partition has been suggested by Boyer into the "fixed"

(liquid-like) and "true" liquid states. The liquid-like transition a", involving motion of the entire molecule, has been observed at T ^

= (1 . 20 - 0.05) T . The liquid-liquid transition T ^ ,

involving segmental mobility in the liquid state, has been usually observed about 40 K above T . g

Thus, it is claimed that the

transition could correspond to the liquid-liquid transition. The liquid-like transition increases with the crosslinking density (or absorbed radiation dose). The ratio a '/a varies in the range of 1.10 to 1.27, and it increases with degree of curing

(Fig.6).

Although the p-transition was clearly observed in most samples, no general trend of the dependence of T^ on the degree of curing could be established. However, within the a' transition T P increases with the curing degree. Very high dielectric dispersion of p-transition observed at the lowest frequency in TSD may be

H201+P25

250

' W/3)

200

• • A 150

6.25

125

kGy 4 6 8 vol %

Fig.6. Transition temperatures by poling current.

369 attributed

to the displacement of free carriers which could be

trapped at the boundaries between different phases. The model mechanism could be proposed: The electronic of the hydrogen bond

following

delocalization

(-C=0-••-H-0-) would allow electrons to jump

from one chain to another. The hydrogen bonds which stabilize the structure in the highly entangled state may thus allow jumping of the delocalized electrons. The a -mechanism may involve

sequences

containing carboxyl groups which are able to be fixed by hydrogen bonds or to fix water molecules. There are many possibilities of bonding between carbonyl groups, either intra- or interchain. So, two neighboring carbonyl groups of one chain may be linked by two water molecules. The kinetic unit would then become larger. The a process can be attributed to the breaking and reforming

hydrogen

bonds with the macromolecules. The rotation of COO dipoles may involve a larger movement of the main chain and could be responsible for 6 process. The y process corresponds to the rotation of the

terminal COOH dipoles. As UP resin exhibits an aggregate

structure, the filler could not penetrate into the aggregates but would fill-up the inter-aggregate has a possibility

inter-bundle space. The aggregate

to undergo thermal motion as a whole unit. Thus,

the a' peak which appears in the TSD spectrum is a liquid-like-type transition corresponding most probably to the thermally

activated

destruction of the aggregate structure. Adding of glass micronodules which are initially completely soluble in the resin resulted in the beads separation, leading to voids and beads. The relatively

separated

small filler beads, situated between the

chains give rise to weak random potentials. Our main assumption is that beads destroy structural ordering of the conducting

chains

which are stabilized by hydrogen bonds. Early treatments of the effect of the filler low concentrations, based on the free volume theory, predict an increase in the transition temperature. Recent result, the depression of the transition temperature,

suggests

that the transition is smeared and the long range order is lost. In the limit of the intermediate concentrations of filler, the off chains correlations are removed and the system consists of decoupled chains with strong on-chain correlation effects. Also, it is expected that the conductivity becomes non-ohmic for large electric field strength. The material with a high

concentration

of beads may be represented as a random mixture with no correlation.

370 It s h o u l d be giving

rise

positions

content.

essential

The

changes

linked UP resin.

the

contributes

in p a s s i n g

manifested The

to the r u p t u r e in the

number

particles,

i.e.

The

to the

chains

of polymer

a reduced density

It c a n be s e e n of the

sition

shifts

beads.

The

temperature

layers, of

due

analogous lers,

to the m e c h a n i c a l

can be physically

a fine

lying

the cavity

surface.

consequently of the m u t u a l composite It m e a n s

This

The

An

interaction

that

in f i l l e d of r a d i c a l s

of

systems

12.5 a n d

in

with par-

the

a-tran-

25 v o l . %

of the

increases

"interfacial

of

surface

structure with is

solid

fil-

polarization"

can be considered

acquire

cavity

of

induction

phenomenon

charges

of the m a t r i x (observed

s t a g e s of

(2 1 2 . 5 v o l . % )

the

in

the

on

is Fig.6)

beads

hardening.

filler

increases

as

the

in UP r e s i n - g l a s s

simultaneously

in

low-content

by m e a n s o f

filler

the

solid

in the n o d u l a r

the conductivity

inversion

of

is a s s o c i a t e d

surface

current

an

filler

lower-temperatures an increase

and

of

of the

This effect, which

of c o m p o n e n t s

and

formation

(6 a n d 8 k G y ) ,

materials was revealed at various

the m o b i l i t y

filler,

surface

in a spherical

particle will

Therefore,

increased.

shifts

for

reinforcement

explained by

particle

lattice, is

the

in the t h i c k n e s s

filler.

effect.

dielectric matrix.

o n the

by

increase

transition.

indicates

resulting

or M a x w e l l - W a g n e r - S i l l a r s spherical

in the

for u n f i l l e d and

the p o l i n g

f r a c t i o n of

temperatures

to the

interface

the

liquid-like

toward

temperatures

increase

As a rule,

volume

This

dependence

indicates

while

that, w i t h the reduction

to curing m e c h a n i s m

the network.

increasing

of

cross-

temperature

bonds

the

layer by c u r i n g

toward higher

filled UP resin

of

shifts

packing

(Fig.4)

surface

physical

no

measurements)

in the e f f i c i e n t d e n s i t y

in the

of p o l y m e r

to f i l l e d

current

bonds w i t h the

(4 k G y )

mobility

the

elements whole mobility

introduction

filler.

the

spectra undergoes

The

which testifies

increase

a-transition

because

independent of

toward higher

region

of p h y s i c a l

the

thickness

some

motion

lower temperatures,

(at 4 k G y ) .

of

upon the

w i t h the c o n t e n t of

ticles.

shifts

structural

temperature,

additional network.

toward

temperature

a'-transition,

toward higher

(in t h e p o l i n g

shifts

a'-process

larger

unchanged,

from unfilled

over

filler content

thus releasing

remains

of d e p o l a r i z a t i o n

It is s h o w n

high-temperature

of m o l e c u l a r

of T S D p e a k s a r e a l m o s t

pattern

that the a - t r a n s i t i o n increasing

that the character

to the d e p o l a r i z a t i o n

temperature filler

stressed

limits

371

achievable crosslinking density. It might seem that since the change in transition temperature is associated with the restriction in molecular mobility, the filler content would have to manifest itself the more, the thinner the polymer layer between two particles of the filler, i.e. the greater its concentration. This is indeed observed in Figure 6 after a minimum on the curve of transition temperature vs. filler content. Thus, the descending part of these curves should be affected not by the thickness of the layer, but by the portion of macromolecules close to the surface, where the mobility of macromolecules is considerably reduced. The minimum in Figure 6 is observed due to an increase in mobility of polymer chains in the interface and is associated with a reduced density of polymer packing on the surface of solid particles. It becomes clear that most of the functional groups on the surface are not directly bonded to the filler surface, due to aggregation present in the UP resin at low filler content.

Conclusions All the preceding experimental facts conclusively show that mixing filler and UP resin has little influence on the intrinsic properties of UP resin in its glass transition range and involves the formation of an interfacial polarization at the phase boundaries. The reasons for the seeming failure of TSD measurements to reveal much about morphology could be: first, the weight fraction of beads was very small; second, curing has been demonstrated to have a strong effect on TSD spectra, but (apparently) only a small effect on morphology; third, considerable MWS (Maxwell-WagnerSillars) effect could be expected to strongly determine measured current values and masks suitable changes in spectra attributable to morphology and interface layer effects.

CLASS

BEADS FILLED EPOXY SYSTEM

INTRODUCED Y.G.

ELASTOMER

: THE

TOUGHENING

EFFECT OF

AN

INTERPHASE

L i n , J . P . P a s c a u l t a n d H.

Sautereau

L a b o r a t o i r e d e s M a t é r i a u x Macromoléculaires - U A C N R S n ° 507 I n s t i t u t National d e s S c i e n c e s A p p l i q u é e s de L y o n - Bâtiment 403 20, A v e n u e A l b e r t E i n s t e i n , 69621 V i l l e u r b a n n e C e d e x , F r a n c e

I ntroduction T h e s t r e n g t h p r o p e r t i e s o f filled e p o x y n e t w o r k s d e p e n d o n many

parameters

s u c h as volume f r a c t i o n ( 1 - 7 ) , size ( 2 - 6 ) a n d s u r f a c e treatments ( 7 , 8 ) o f the f i l l e r s , a n d chemical s t r u c t u r e a n d c r o s s l i n k i n g d e g r e e o f the m a t r i x . T o i n c r e a s e the s t r e n g t h p r o p e r t i e s o f s u c h materials, two methods are c u r r e n tly u s e d : (i )

the reinforcement o f the matrix itself d u e to the i n c l u s i o n o f r u b b e r

particles ( 9 ) , u s u a l l y C T B N

(10,11).

(it) the improvement o f interfacial s t r e n g t h b y means o f c o u p l i n g a g e n t s (6,8,12). H o w e v e r , it is w e l l - k n o w n that failure almost a l w a y s o c c u r s from m i c r o - d e f e c t s orinhomogeneiities in the materials, w h i c h lead to localised s t r e s s e s

considera-

b l y b e y o n d t h e a v e r a g e s t r e s s in the b u l k . I n the c a s e o f g l a s s filled e p o x y s y s t e m s the e x c e s s s h r i n k i n g o f the matrix c a u s e s a h y d r o s t a t i c c o m p r e s s i v e stress. H e r e , we s t u d y the t o u g h e n i n g effect o f an elastomer l a y e r on the g l a s s b e a d s , t h i s i n t e r p h a s e may r e d u c e the localised s t r e s s e s o n the filler s u r f a c e . few w o r k s h a v e been r e p o r t e d o n s u c h a s u b j e c t ( 1 3 , 1 4 , 1 5 ) .

Only

Untreated and

si lane treated g l a s s beads are also u s e d in t h i s w o r k for c o m p a r i s o n .

Experimental M a t e r i a l s . T h e chemical p r o d u c t s u s e d in t h i s s t u d y are listed in T a b l e 1. The glass beads ( S O V 1 T E C

50 A O ) , in the r a n g e o f 4-44 um, h a v e an a v e r a g e

diameter o f 26 pm. P r e p a r a t i o n o f the a d d u c t . T h e p r e p o l y m e r D G E B A a n d the C T B N are m i x e d with a c a r b o x y l - t o - e p o x y

ratio

r 1 = 0.5. T h e reaction is c a r r i e d out at 85°C,

u n d e r mechanical s t i r r i n g a n d n i t r o g e n flow, u s i n g 0.15 % of t r i p h y l p h o s p h i -

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

374 ne as catalyst. T h e mechanism of this reaction has been already d i s c u s s e d (10.11). The analyses b y C P C show that the average molar weight of the C T B N D C E B A mixture increases linearly with time for 17 h o u r s . After 19 h o u r s the reaction is ended and then a liquid diamine ( M N D A ) is added whose primary amine h y d r o g e n atoms can react slowly at room temperature. A ratio of amineto-epoxy r 2 = 2 is chosen, assuming that half of epoxy g r o u p s has already reacted in the early stage. T h i s mixture is kept at 22°C for 4 d a y s ; the C P C analysis shows that all the residual D G E B A molecules have disappeared and that a further increase in average molar weight has

occurred.

Surface treatment o f g l a s s beads. Two surface treatments, with respectively adduct and silane coupling agent A187, are adopted. In the first case, 100 grams of g l a s s beads are mixed with 175 ml of an adduct-methyl ethyl ketone solution (2 %), at 22°C for 10 minutes. Then the solvent is evaporated u n d e r vacuum, and the beads are dried u n d e r vacuum at 120°C. T h e coated adduct quantity is 3.5 grams for 100 grams of glass beads. In the second case, silane coupling agent is used as indicated by the manufacturer. After both treatments the beads are passed t h r o u g h t a sifter of mesh size 200 pm.

Table 1. Chemical

products

Materials

Epoxy prepolymer ST = 380 g/aol

Chemical formula

Trade mark

CH2-CH-CH2 4 - O - ^ Y C - / o V o - C H 2 - C H - C H 2 4 X

O

L

CH«

OH

0-/OVC-A\-CH2-CH-CH2 0.2

3

CH,

O

DGEBA 0164 Bakelite

3

dijjlycidyl-ether of blghenol A Hardener DDA

NH2 -

C s N ' C a N ' NH 2

Catalyst BDMA

Rubber CTBN

( C ^

CHJ - N

HOOC

-

Dicyan^Uamine

VE 2560 Bakelite

- t CHJ )2

Benzyl dimethyl amine

(CH2-CHCH-CH2)5 —(CH2-CH) - -

ST" = 3500 g/mol

COOH

Carboxyl terminated butadiene acrylonitrile

CN

Silane

C H 2 - CH - CH Z - O - ( O i ^ — ^

SI (OCHJ> 3

Hycar C T B N 1300 x 8 Goodrich

10

Coupling agent

Fluka

( Y -glycidoxytrlmetboxysilane)

A187 Union Carbide

o"' CH,

Liquid diamine (MNDA)

1

NH, - C - \ 2 I \ CHJ

R

V

M

!

A. / CH, 3

1-0 p.menthane diamine

Aldrich

Preparation of composites. T h e matrix is formed b y D C E B A and D D A , with an amine-to-epoxy ratio r 3 = 0.6, 1 p h r of B D M A is used as accelerator. The glass beads are added according to different volume fractions (10 %, 20 %, 30 %). The materials are mixed by mechanical s t i r r i n g u n d e r vacuum, at 60°C

375 f o r 1 h o u r . T h e n t h e m i x t u r e is c a s t e d i n t o a P T F E - c o a t e d mould a n d c u r e d at 160°C f o r 1 h o u r followed b y 1 h o u r at 180°C. Instrumental.

T h e s y n t h e s i s o f a d d u c t is m o n i t o r e d b y gel p e r m e a t i o n

(CPC)

w i t h a WATERS a p p a r a t u s , t e t r a h y d r o f u r a n e is u s e d as s o l v e n t . T h e m e c h a n i cal p r o p e r t i e s a r e measured b y t h r e e - p o i n t s b e n d i n g t e s t o n a t e n s i l e machine (DY 14 A D A M E L LHOMARGY) a n d b y impact t e s t w i t h an i n s t r u m e n t e d

set-up

( 1 7 ) . S c a n n i n g e l e c t r o n m i c r o g r a p h s (SEM) o n f r a c t u r e s u r f a c e s a r e o b t a i n e d w i t h a J E O L T200 a p p a r a t u s .

Dynamical p r o p e r t i e s are r e c o r d e d w i t h a m i c r o -

torsional set-up developped by J . Y .

CAVA ILLE

(18).

Results a n d D i s c u s s i o n

Since t h e r e a c t i o n mechanism o f t h e D C E B A - D D A s y s t e m is r a t h e r complex

(19,

20), the final p r o p e r t i e s o f the n e t w o r k formed depend s t r o n g l y on t h e cure conditions (21).

In this w o r k , the present c u r e conditions are chosen in o r d e r

to g e t a m a t r i x h a v i n g h i g h e r s t r e n g t h t h a n t h e m a t r i x - f i l l e r i n t e r p h a s e , t h e i n t e r p h a s e e f f e c t w i l l be emphasized.

then

DSC measurements show t h a t , i n t h e -

se c o n d i t i o n s , t h e chemical r e a c t i o n is m o s t l y e n d e d ( n o e x o t h e r m a l e f f e c t ) t h e m a t r i x t h u s f o r m e d has a glass t r a n s i t i o n t e m p e r a t u r e T g at 127°C, pendent o f the filler surface treatment and the filler

E v i d e n c e f o r i n t r o d u c e d elastomer i n t e r p h a s e .

and

inde-

content.

T h e amount o f a d d u c t in c o m p o s i -

tes is v e r y small, e . g . o n l y 1 . 2 % i n composite f i l l e d w i t h 20 % o f g l a s s b e a d s . So i t c a n n o t

be e v i d e n c e d b y DSC m e a s u r e m e n t s .

On t h e o t h e r h a n d ,

becau-

se a d d u c t is coated o n t h e f i l l e r s u r f a c e , an elastomer i n t e r p h a s e is f o r m e d b e t w e e n t h e f i l l e r a n d m a t r i x . S u c h a c h a n g e in i n t e r p h a s e p r o p e r t i e s c a n be r e v e a l e d b y d y n a m i c a l p r o p e r t i e s m e a s u r e m e n t s , as i n d i c a t e d b y L I P A T O V

and

coworkers (15,16,17).

in-

I n t h e p r e s e n t w o r k , t h e p r e s e n c e o f an elastomer

t e r p h a s e is e v i d e n c e d b y d y n a m i c a l measurements w i t h a t o r s i o n f r e q u e n c y o f 0.01 Hz

: beside t h e @ r e l a x a t i o n peak o f t h e e p o x y m a t r i x at 180 K , a small

r e l a x a t i o n peak a p p e a r s at 217 K ( F i g .

1 ) , w h i c h c o r r e s p o n d s to t h e g l a s s

s i t i o n o f t h e elastomer i n t h e i n t e r p h a s e .

I t is also o b s e r v e d t h a t t h e p r e s e n c e

o f elastomer i n t e r p h a s e s l i g h t l y s h i f t s t h e 0 r e l a x a t i o n peak o f t h e m a t r i x to lower

temperatures.

tran-

376

Fig.

1.

tg & versus

T for

20 % filled

composites. (h) adduct Torsion

treated

frequency

T o u g h e n i n g effect d u e to elastomer i n t e r p h a s e .

; (*)

untreated

0. 01 Hz

T h i s t o u g h e n i n g effect i s well

e v i d e n c e d b y both t h r e e - p o i n t s b e n d i n g test a n d impact test at room t e m p e r a t u r e , e s p e c i a l l y w h e n the filler f r a c t i o n r i s e s u p to 30 % ( F i g .

2). It is o b s e r -

v e d that the f l e x u r a l m o d u l u s d o e s not v a r y with s u r f a c e treatment,

proving

that at low s t r a i n s n o s i g n i f i c a n t influence o f the interface c a n be e v i d e n c e d (3,4).

Fig. (a) (

(b)

2. Mechanical

three

points

) untreated

tests

bending ; (-.-.-.)

for

; (b)

30 % filled

Charpy

silane

composites,

instrumented

treated

(

impact. ) adduct

treated

O n the c o n t r a r y , the s u r f a c e treatments h a v e a s t r o n g e r i n f l u e n c e o n t h e s t r e n g t h p r o p e r t i e s o f the composites. T h e composites filled with treated e x h i b i t a l w a y s h i g h e r s t r e n g t h t h a n t h o s e with u n t r e a t e d b e a d s ( F i g . t i n g an improvement in i n t e r p h a s e p r o p e r t i e s .

beads

3),

indica-

In a d d i t i o n , it seems that the

a d d u c t c o a t i n g has still better enhancement effect t h a n the silane treatment. In Fig.

3, we note that the composites filled with a d d u c t c o a t i n g b e a d s

have

the h i g h e s t resilience ( w o r k to b r e a k / c r o s s section area) ; it may be d u e to

377 the e n e r g y a b s o r p t i o n o f the elastomer

Fig.

3. Resilience

fraction. (0)

versus

filler

(

) impact

(

) three-points

untreated

beads

layer.

Fig.

;

; (o)

versus bending

silane

treated

4. Initiation filler

energy

fraction

in

fraction impact

behaviour beads

; (A ) adduct

treated

I n impact t e s t , the u s e of an i n s t r u m e n t e d s e t - u p allows u s to a n a l y s e t h e a b s o r b e d e n e r g y . A s w i t h A D A M S a n d WU (22) the total e n e r g y may be s e p a rated into two p a r t s

:

( i ) t h e e n e r g y o f initiation ( w o r k d o n e to r e a c h the g r e a t e s t a p p l i e d stress)

;

( i i ) the e n e r g y o f p r o p a g a t i o n . T h e balance between t h e s e two p a r t s g i v e s u s information about t h e mechanism o f deformation a n d the a b s o r p t i o n o f e n e r g y d u r i n g impact test. We see c l e a r ly o n F i g .

4 that the p r e s e n c e of the a d d u c t k e e p s n e a r l y c o n s t a n t t h e v a l u e

o f Ej^ when the filler f r a c t i o n i n c r e a s e s . A similar effect i s g e n e r a l l y o b s e r v e d in r u b b e r t o u g h e n e d e p o x y

system.

F r a c t u r e s u r f a c e m i c r o s c o p y . O n the m i c r o g r a p h s p r e s e n t e d in F i g . 5 we o b s e r ved clearly : (i) poor c o h e s i o n between u n t r e a t e d g l a s s filler a n d m a t r i x , the e x p o s e d s u r f a c e o f filler is g l o s s y ( F i g .

5a) ;

(ii) g o o d c o h e s i o n between silane treated g l a s s filler a n d m a t r i x , t h e r e are b r o k e n polymer pieces c o v e r i n g the filler s u r f a c e ( F i g . 5b)

;

(iii) trace o f p l a s t i c deformation o f polymer materials o n the filler s u r f a c e (Fig.

5c).

The last observation proves the e n e r g y absorption effect of the introduced elastomer interphase.

a)

untreated

b) silane

treated

c) adduct Fig.

5. Scanning

surfaces

treated

electron

on composites

microphotographs

filled

with glass

on the beads.

fracture

379 Conclusion

I n t h e p r e s e n t s t u d y , we have e x h i b i t e d a s p e c i f i c r e l a x a t i o n peak o f t h e a d d u c t i n t e r p h a s e coated on g l a s s beads. T h e mechanical p r o p e r t i e s show t h a t s u c h a c o a t i n g has b e t t e r r e i n f o r c e m e n t e f f e c t t h a n t h e u s u a l silane t r e a t m e n t , e s p e c i a l l y i n impact b e h a v i o u r .

T h e s e f i r s t e x p e r i m e n t s may c e r t a i n l y be i m -

p r o v e d w i t h o t h e r compositions o f t h e elastomer a d d u c t , a n d a p p l i e d t o f i b r e reinforced composites.

Aknowledgement

T h e a u t h o r s a r e g r a t e f u l l to D r . J . Y . C A V A I L L E a n d C . J O U R D A N f o r help in dynamical

their

measurements.

References 1. M a n s o n , J . A . , S p e r l i n g , L . H . Plenum P r e s s , New Y o r k .

1976. I n

2. S p a n o u d a k i s , J . , Y o u n g , R . J .

1984. M a t . Sei. 1j), 173.

3. S p a n o u d a k i s , J . ,

1984. M a t . S e i .

4. S a h u , S . ,

Young, R.J.

Broutman, L.J.

: Polymer B l e n d s a n d C o m p o s i t e s .

19, 487.

1972. Polym. E n g . S e i . , 12, 2,

91.

5. M o l o n e y , A . C . ,

Kausch, H . H . , Stieger, H.R.

1984.

J. Mat. S e i . ,

6. M o l o n e y , A . C . ,

Kausch, H . H . , Stieger, H . R .

1984. J . M a t . S e i . ,

18, 208. 19, 1125.

7. I s h i d a , H . , K u m a r , C . 1983. I n I n t e r f a c e s . Plenum p r e s s , N . Y .

: Molecular C h a r a c t e r i s a t i o n o f Composites

8. Morel I , S . H .

Proc. A p p l . ,

1981. P l a s t . R u b b .

9. B u c k n a l l , C . B . 10. B a r t l e t , P h . , 11. R i e w , C . K . ,

13. A b a t e , C . F . ,

1977. I n : T o u g h n e d p l a s t i c s . A p p l . Sei. P u b l .

Pascault, J . P . ,

Sautereau, H. J. Appl.

R o w e , E . M . , S i e b e r t , A . R.

12. B r o u t m a n , L . J . ,

179.

Sahu, S.

Heikens, D.

Polym. Sei. ( i n

1976. E d . A d v .

1971. M a t . S e i .

Eng.

98.

1983, Polym. Com. 24, 137.

14. E a s t m o n d , C . C . , M u c c i a r e l l o , G .

1982. Polymer 23,

London.

164.

press).

Chem. S e r i e s ,

326.

380 15. Lipatov, Y u . S . ( R o s v i z k y , V . E . , S h i f r i n , V . V . 1982. J. A p p l . Polym. Sei. 27, H55. 16. Lipatov, Y u . S . # 1977 : I n Physical Chemistry of Filled Polymers, Khimiya Moscow. 17. B a b i c h , V . F . , Lipatov, Y u . S . 1982, J. A p p l . Sei. 27, 53. 18. Etienne, S . , Cavaill£, J . Y . , Perez, J . , Point, R . , Salvia, M . 1982. R e v . Sei. Instrum. 53, 1261. 19. Galy, J . , 1985, P H D T h e s i s ,

Lyon.

20. Zahir, S . A . 1980. T h i r d International Conference in O r g . Coat. S e i . , Athens. 21. Lin, Y . C . , Sautereau, H . , Pascault, J . P . ( i n p r e s s ) . 22. Adams, C . C . , Wu, T . K . 1981, S P E A N T E C ,

185.

STRUCTURE AND MECHANICAL PROPERTIES OF POLYMER-PHENOLIC MICROSPHERE COMPOSITES D. ¿uchowska, 0. Malczewski, L. Woiniak Institute of Organic and Polymer Technology, Technical University of Wroclaw, 50-370 Wroclaw, Poland

Introduction The polymer composites obtained by mixing a polymer with microspheres acting as a filler are sometimes called the syntactic plastics. They have specific polymer-polymer-void heterogeneous structure and their properties depend on many factors, the most important of which are the size of microsphere grains, method of their admission to the polymer continuous phase, and adhesion on the polymer-microsphere interphase. The properties of composites containing glass beads or microspheres have been reported in [1-6], while, more generally, the syntactic plastics are discussed in a monograph [7], The aim of this work was to characterize the morphology and properties of polymer composites containing various amounts of phenol-formaldehyde microspheres distributed in isotactic polypropylene (ductile) and in crosslinked epoxy resin (brittle material). The mechanical stability of the phenol-formaldehyde microspheres, especially against hydrostatic pressure, is much lower than that of glass microspheres [7]. When subjected to shear forces during homogenization and moulding, these microspheres can be seriously affected, too. However, in order to determine the economical value of the phenolic microspheres, the typical processing conditions were used for preparing the composites.

Experimental Materials The phenol-formaldehyde microspheres (PFM) were obtained on a laboratory scale from a commercial resol resin FK-74. The PFM s

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed iri Germany

382 had density 0.123 g/cm 3 and grain diameter 80-300 urn. They contained 92% of acetone insoluble part, and the weight losses of 10, 20, and 50% were at 130, 230, and 380 °C, respectively. Isotactic polypropylene (iPP) Malen 3 400 (Poland) having flow index 3 g/min,

and epoxy resin Epidian 5 (Poland) of viscosity of ft

80 000 cP at 20 C and density of 1.17 g/cm were used. The epoxy resin (EP) was crosslinked with triethylenetetramine (TETA). Samples iPP - PFM Composites Polypropylene melt was mixed with PFM's at 240 °C in a screw extruder, followed by granulation. The test samples were formed by injection moulding under pressure of 15 MPa. The mould temperature was 60 °C. EP - PFM Composites The samples were formed using a pressureless method. The epoxy resin, 90 wt.%, was mixed with PFM's and 10 wt.% of TETA. The mixture was degased under reduced pressure. The volume fraction of microspheres was calculated in relation to the total volume of the resin and TETA. The setting was carried out at 20 °C for 24 h and at 80 °C for 2 h.

Procedures The mechanical properties were measured using a tensile strength machine of the Instron type. The rates of deformation was 50 mm/min for iPP-PFM, and 10 mm/min for EP-PFM composites. The standard deviations of the stress measured for 7 samples were 0.24 MPa for iPP-PFM, and 2.1 MPa for EP-PFM composites. The thermal stability was measured on a Oerivatograph (Hungary) at the rate of temperature rise 2.5 K/min . The physical structure was estimated by examination of the surfaces of samples and

their fractures by using a scanning

electron microscope (SEM), Cambridge Stereoscan (England). The surface were coated with gold.

383 Results and Discussion The PFM's had a spherical shape and were filled with air. The walls had a thickness of ca. 5 jjm, as determined by SEM examinations. A fraction of PFM's (ca. 20 %) was found to have shell walls broken. Due to a strong electrostatic charge the PFM's tended to form agglomerates. An addition of even a small amount of the phenolic microspheres to iPP yields a rapid decrease in the elongation at break, of this polymer (Table 1). This is quite general that thermoplastic materials respond in this way when filled with hard spherical particles. The elongation above the yield point, £ y , slightly increases at the volume fraction of microspheres 0.01, thus indicating an affect of toughening. Table 1. Composition and Properties of iPP-PFM Composites Volume fraction of PFM

Density / 3 g/cm

Yield stress MP a

0

0.88

0.005 0.010

0.87 0.85

31.7 32.8 33.0

0.100

0.86

31.7

Elongation, % Specific work, 3/cm W £y £b y W b 17

447

17 19 12

178 34 13

4.37 4.01 4.19 2.97

3

89.34 25.86 7.64 2.97

For the brittle epoxy based composites, a reduction in stress accompanied by a slight decrease in elongation is observed at the content of PFM's exceeding 0.1 (Table 2). The interaction on the polymer-microsphere interface have been determined from the stress-strain relationship. According to the assumption of Nicolais and Narkis [1],if there were no adhesion between the filler and polymer,then the total stress would be transferred by the continuous polymer phase alone. The geometric considerations lead to the following relationship between the yield stress in composite, (5 c , the yield stress within polymer matrix, 6 and the volume fraction of filler, V^ : 3 6 cy - 6 my (1 - 1 .21 vf/ ) f

(1)

384

Table 2. Volume fraction of PFM

Composition and Properties of EP-PFM Composites Yield Elon- Spectatress gation fie

Density g/cm3

M P a

expl. cald.

3

Strength compression MPa

bending MPa

impact

121

92

23.2

*

0/cîn

32.99

4.7

0.57

32.31

4.6

0.56

93

52

22.95

4.1

0.35

75

41

O/m2

0.2

1 .22 1 .22 1 .07 1 .11 0.97 1 .00

0.3

0.86

0.89

56

33

7.8

0.4

0.81

0.78

51

32

6.4

0.5

0.72

0.67

23

5.8

0

0.1

10.2 8.1

For a filler with particles that do not deform upon stress, the deformation up to the yield point,

&c

, can be calculated

using the relationship proposed by Nieison [8]: £

C

y

-

V

1

• u H0 n 0 "O > 4-* •ri e 4-* * 0 N O a i-i C M a. 3 E O • o r-H CT u (S 0 O E •O « O . 1 t. o a. rH 0 i f a. c a U) 0 H 3 "O o O c .—. — r 1 • •— -O 0 O Œ er u) U 0 s ID C u. •ri a. X (D iH i «0 a. 0 X» a. 3U • r i •H (D 4-» < H 0 :. O u 0 o 4- 0 •o iH ü « 0 0 C •ri >«0 • r i U L. O 0 CL «4- co •C o a i-l • 0 0 0 c o 4» o o L. C c •H U 0 0 4-» • r i s T> 0 E • r i 1. C e 0 u. O 0 a. o -ri a rH X 0 •o 0 O 0 T3 c 0 0 0 -C 4- JC .c 1- o a 1-

«

IedW]

kQ

£ on

..

o> •ri LL

386

Fig.2.

SEM microphotographs of iPP-PFM composite

surface.

The content of P F M ' s : 1% by volume.

Fig.3.

SEM microphotographs of the surface of fracture by stretching a sample of iPP

(a) and iPP-PFM

obtained

compo-

site containing 1% of the microspheres, by volume,

(b).

387 M i c r o c r a c k s s p r e a d i n g over the m a t r i x a l o n e that b y - p a s s s p h e r e s can a l s o o b s e r v e d

Fig.4.

micro-

(Fig.4).

S E M m i c r o p h o t o g r a p h of the s u r f a c e of f r a c t u r e of EP-PFM composites PFM vol. fraction 0.1, sample by

broken

bending.

The thermal s t a b i l i t y of i P P - P F M c o m p o s i t e is h i g h e r than of the m a t h e r p o l y m e r . T h e t e m p e r a t u r e at w h i c h

that

the w e i g h t

loss

r e a c h e s 50 w t . % m o v e s t o w a r d s h i g h e r t e m p e r a t u r e . No such an e f f e c t w a s o b s e r v e d for E P - P F M

composites.

On the o t h e r h a n d , the E P - P F M c o m p o s i t e s , u n l i k e the o n e s , have d e n s i t i e s equal

approximately

assuming volume additivity

(Table

The r e l a t i v e s p e c i f i c per d e n s i t y u n i t s ,

strength

decreases

to that c a l c u l a t e d

with increasing volume

tensile

>

calculated

fraction

the p a r t i c u l a r s t r e n g t h s to

i n m i c r o s p h e r e content d e c r e a s e s in the c o m p r e s s i o n >• b e n d i n g

The b i g g e s t c h a n g e s in

>

(

bending,

0.1).

volume

Further increase

v o l u m e f r a c t i o n of the m i c r o s p h e r e s in the e p o x y resin up to 0 . 5 b r o u g h t a b o u t a s m a l l e r d e c r e a s e ties .

the

impact

strength during compression,

phenolic microspheres

of

order

and i m p a c t tests w e r e o b s e r v e d for s a m p l e s w i t h a low f r a c t i o n of

by

2).

of E P - P F M c o m p o s i t e s ,

m i c r o s p h e r e s . T h e s e n s i t i v i t y of increase

iPP-PFM

of

those

in

matrix proper-

388 Conclusions The phenolic microspheres reinforce polypropylene when their volume fraction is 0.01. For crosslinked epoxy resin, the toughening is observed at the volume fraction of microspheres of

0.1.

As the volume fractions of microspheres increase above these values, the specific work of deformation, decreases due to - reduction in plasticity (iPP-PFM composites) or . reduction in stress (EP-PFM composites) One can expect an adhesion between the phenolic microspheres and polypropylene or epoxy resin to exist, although other methods of study the composites are required to confirm this conclusion.

References 1.

Nicolais, L.,

N.Narkis:

1971. Polymer Eng. Sei.

2.

Nicolais, L., 15, 35.

O.Acierno,

3.

Nicolais, L.,

L.Nicodemo. 1973 . Polym.Eng.Sci. ^3, 469.

4.

Okuno, K.,

5.

Nicolais, L.

6.

Pegoraro, M., A. Penati, E. Cammarata, M.Aliverti. 1984. In: Polymer Blends, Processing, Morphology, and Properties V 2 (M. Kryszewski, A. Gal^ski, and E. Martuscelli, eds.). Plenum Press. New York and London, p. 205.

7.

Berlin, A.A., F.A. Szutow. 1980. Reinforced gas-filled plastics (in Russian), Moscow, Chimia, p. 158.

8.

Nielsen, L.E.

O.Oanacek.

194.

1975. Polym.Eng.Sci.

R.T. Woodhams. 1974. 0.Cellular Plastics 10.237. 1975. Polymer Eng.Sei. 1J5, 137.

1966. O.Appl.Polym.Sci. 1_0, 97.

CATALYTIC EFFECT OF THE SOLID ACIDS AT AMINO RESINS SOLIDIFICATION

V.M. Tcheshkov, M.M. Natova, G.Z. Zachariev Central Laboratory of Physico-Chemical Mechanics of the Bulgarian Academy of Sciences, Sofia 1113, Bulgaria G.V. Kozlov, T.M. Morozova Institute on Physical Chemistry of the Academy of Sciences of the USSR, Moscow 117312, USSR

Introduction As is known, cross-linking of amino resins and urea-formaldehyde)

(melamine-formaldehyde

occurs in solutions under the action of or-

ganic and inorganic protonic acids which are homogeneous of the process

catalysts

(1). The mechanism of reaction is characterized by

formation of intermediate products - carbonium ions, obtained by proton adherence to the oligomer methylol groups, and

then

by

interaction between these ions and methylol or amide groups. What is obtained as a result is a polymer space network while the catalyst nature affects the ratio between methylene and methylenether links, potential number of cross-links, as well as size and size distribution of supermolecular formations. The above

characteris-

tics affect strongly polymer crack resistance. Resins which solidify under the action of the catalysts considered are not sufficiently crack resistant. Partial elimination of this defect can be attained by introducing a number of inorganic fillers. Shortage of crack resistance can be entirely eliminated by filling the oligomer with sulphates of alkali-earth metals, e.g. BaSO^ chemical activity of these sulphates

(2). However,

manifests itself during cu-

ring acceleration as well. There are various explanations of the effect of the filler, e.g. existence of a difference between the surface energy of the oligomer and that of the inorganic

substance,

catalyst predominant adsorption on the filler surface, orientation effects, steric hindrances, presence of polar groups on the surface, etc. (3).

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

390

We suppose that decisive part is played on behalf of the acidic nature of alkali-earth metal sulphates, while the

process as a

whole can be explained on the basis of homogeneous-heterogeneous acido-basic catalysis. Thus, ionization can be stimulated and ions can participate in adsorption interaction with active acidic sites on the surface. However, insufficient attention is paid to these sites when studying polymerization and polycondensation,

although

that they can affect conditions of polymer formation. Direct study of amino resin adsorption, however, is slowed down due to a number of experimental difficulties, i.e. limited water

solubility

of the oligomer, extremely high polydispersity, etc. This is the reason why acrylamide is used as a model compound in this study. It has the same functional groups - carbonyle and amide ones, as these in amino resins. The present paper attempts to establish filler acidity - catalytic activity and selectivity

correlation

during amino resins curing.

Experimental Part We use the following inorganic

substances to study

adsorption:

silica, titanium oxide, aluminium oxide, calcium phosphate, barium sulphate and calcium fluoride. These adsorbents have been used as fillers in previous studies basis of acido-basic

(2). Their acidity is evaluated on the

interaction as presented by BrfSnsted and

Lewis and is estimated by butylamine titration (A). Following Br^nsted and Lewis, solid acid is considered to be any substance which is capable of donating a proton or of accepting an electron pair, while solid base - any substance which is acceptor of protons or donor of electrons on the solid surface

(5). Moreover, strength of acidic sites

(P K a ) is determined by the

surface

capabili-

ty to transform base molecules from neutral into conjugated acidic form. The p K a ~ v a l u e s of the most popular indicators are within +6.8 i -8.2 range which corresponds to sulphuric acid concentra_g tion of 8x10 -f 90 wt.°&. Moreover, the smaller the pK a -valuesthe higher the strength of the surface acidic sites. We have not included in this study solid acids of medium and high strength (silica-alumina ones ) which do not affect polymerization

practically

(2), as well as water-soluble solid acids. Since

391 water solutions carry out adsorption, adsorbents are net thermally treated to avoid change

of surface acidity

(5). The specific

surface is obtained by nitrogen low-temperature adsorption while acrylamide is obtained by

methanol

recrystallization. The quan-

tity of acrylamide adsorbed at 25°C is determined by employing a weighing

method. Acrylamide infrared spectra are obtained by

using a "Spekord 75 IR" spectrometer.

Results Table 1 presents strength and concentration of acidic sites on the surface of oxides and

salts used in the experiment. SiC^ has

the strongest sites while A ^ O j lacks them at all. Sites concentration in BaSO^, TiC^ and CaF£ is the highest one. Quantities of acrylamide with maximum adsorption on the surface are given as well. Adsorption attains its highest values on CaF^,

while these

values are lower by an order and differ insignificantly from one another, regarding the rest of the adsorbents. Table 2 shows infrared spectra of condensed acrylamide and adsorbed on the inorganic surfaces. Adsorption bands, which are characteristic for valence vibration of carbonyle groups ?(C=0) and for deformation vibration of amide groups cTCNH^), are separated from the spectra. However, these bands are sensitive to adsorption interaction. Spectra of acrylamide adsorbed on SiC^, TiC^, A ^ O ^

and

CaF2 do not differ from spectrum of the initial acrylamide. However, splitting of the N ^ - g r o u p adsorption bands is registered during adsorption of acrylamide on BaSO^ and C a j C P O ^ ^ 1 that, significant shift of the carbonyl cy region of 27 cm ^ occurs in BaSO^

More than

group to the low-frequen-

spectrum.

Discussion Results show (Table 1) that acidity of adsorbent surfaces does not correlate with values of maximum adsorption on the latter, but adsorption values are not always indicative of processes that take place in adsorbate-adsorbent systems. Significantly richer information can be gained from spectroscopic investigations of the ad-

392 Table 1. Acidity of Adsorbents and Maximum Acrylamide Adsorption Adsorbent

Acidic sites

concentration,

H x 1 0 + \ meq/m 2 , for p K & +3.3 Si0 2

+4.8

+6.8

sorption, A, mg/m 2

S.S

-

BaS0 4

-

16.67

-

1.67

Ca 3 (P0 4 ) 2

-

-

0.61

1.46

CaF 2

-

Ti0 2

-

A1,0,

-

-

Maximum ad-

12.5

0.89

-

3.33

17.5

6.66

4.33 2.81

Table 2. Vibration Frequency of Carbonyle and Amide Groups of Acrylamide Wavenumber, Acrylamide cm - 1

(TCNH2) $(C=0)

Acrylamide, adsorbed on Si0 2

1610 1677

1610 1677

BaS0 4 Ca 3 (P0 4 ) 2 CaF 2

[1623

jl 622

\1586

[159 3

1650

1677

Ti0 2

A1203

1610

1610

1610

1677

1677

1677

sorbed substance. In fact, spectrum of the adsorbed acrylamide differ sharply during

transitioni from one inorganic surface to

another (Table 2). Splitting of adsorption bands of NH 2 -groups, recorded for BaS0 4 and Ca 3 (P0 4 ) 2 , can be related to the interaction of their aprotic acidic sites with acrylamide. This takes place by transition of the individed electron pair of the acrylamide nitrogen atom to the surface electron acceptors. So, compounds of coordination character are formed on the surface. Similar effect - splitting of the vibration band 1 0 v o l % lower than the experimental values. The reason for this is that the particles geometry, packing fraction and interaction between polymer and filler are not involved sufficiently in the theory. Because the effective filler concentration is always greater than the analytical one, Nielsen (15,1b) reduced the filler volume scale by introducing a maximum packing fraction of the filler. It is reached at 74 vol% (closest spherical packing) and varies with the particle shape and sort of packing (17). The effective filler fraction is given by Nielsen as:

Y err

' YF

+

3(20 v%): (a) one-step mixing; (b) EPDM mixed with CaCC>3, added to PP.

543 i

i

log G' Pa 9 - o (PP*CaC03)*EP0M e (PP+EPOMl + CaCOj 9 PP+IEPDM + CaCOj) 8 - • PP+EPDM + CaCO^ log G' Pa 7 i -100

I 0

°c

100

Fig. 8. The effect of mixing sequence on the storage and loss shear moduli of the composite PP(60 v%)/EPDM(20 v%)/CaC03(20 v%).

Conclusions Dynamic mechanical measurements and ESM micrographs provide evidence that EPDM elastomer inclusions and surface treated calcium carbonate filler particles are separated in the polypropylene matrix. The rubber and filler phases do not interact and their effects on the stiffness and damping of the composites are apparently independent of each other. The morphology of the ternary composites has been found to be independent of the sequence in which the components are mixed.

References 1. Bucknall, C.B. 1977. Toughened Plastics. Applied Science Publ., London. 2. Kolarik, J., G.L. Agrawal, Z. Krulis, J. Kovar: Polym.Composites (submitted). 3. Bucknall, C.B., C.J. Page. 1982. J.Materials Sci. YT_, 808. 4. Stamhuis, J.E. 1984. Polym.Composites 5, 202.

544 5. Pukanszky, B., F. Tüdös, T. Kelen: Polym.Composites (accepted). 6. Kolarik, J. 1 982. Adv.Polym.Sei. _46, 119. 7. Kerner, E.H. 1956. Proc.Phys.Soc. 69B, 808. 8. Nielsen, L.E. 1974. Mechanical Properties of Polymers and Composites. M.Dekker, New York. 9. Nielsen, L.E. 1978. Predicting the Properties of Mixtures. M.Dekker, New York. 10. Pukanszky, B., J. Kolarik, F. Lednicky , in this book. 11. Ward, I.M. 1971. Mechanical Properties of Polymers. Wiley, London.

MECHANICAL PROPERTIES OF POLYETHYLENE AND POLYPROPYLENE FILLED WITH CALCIUM CARBONATE

J. Kucera Research Institute of Macromolecular Chemistry 656 49 Brno, Czechoslovakia J. Kolarik Institute of Macromolecular Chemistry, Czechoslovak Academy of Sciences, 162 06 Prague 6, Czechoslovakia

Introduction Though thermoplastics filled with inorganic particulate fillers are generally known and widely utilized, their elastic and ultimate properties are difficult to foresee in a great detail, especially in the case of semicrystalline matrices. Our preliminary study on polyolefin composites (injection moulding types filled with calcium carbonate) has shown that mechanical properties of linear polyethylene

(PE) composites

are monotonical functions of

the filler fraction, while the properties of isotactic polypropylene (PP) composites pass through extremes. This communication is concerned with the effects of filling and of the matrix supermolecular structure (controlled by the cooling rate) on yield stress, strain at break, dynamic mechanical behaviour and impact strength of composites.

Experimental Matrices:

Linear PE Liten MB 57 (random copolymer with 1 % of pro-

pylene) and isotactic PP Mosten 58.412 were products of Chemopetrol (Czechoslovakia). Their melt index was 4 g / 10 min (190°C , 21 N ) and 2.6 g / 10 min ( 2 3 0 ° C , 2 1 N ) , Filler:

respectively.

Microground calcium carbonate with calcitic structure,

trade mark Durcal 2 (Omya, Pliss-Staufer AG, Switzerland). About 50 or 97 % of particles were smaller than 3 or 10 ym , respectively.

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

546 The particle surface was treated with 0.3 w% of stearic acid and 0.3 w% of calcium stearate. Sample preparation: Granulate species, 6 of each type, of PE-CaC03 (PE-C) and of PP-CaC03 (PP-C) composites were prepared in a twin screw extruder Werner-Pfleiderer ZDSK 53 VAR . The filler concentration was 0 to 30 v% in both series. Using a simple picture frame form, plates were compression moulded, of which samples were punched or machined for mechanical testing. The moulding temperature was 180°C for PE and 210°C for PP. Differing supermolecular structures were produced by fast (200 K / min , symbol F) or by slow (3 K /min , symbol S) cooling of the plates. Testing: Tensile experiments were performed with an Instron Tester TT-CM (strain rate 300 %/min). Impact testing was carried out by means of a Charpy pendulum Zwick 5102 (impact speed 2.9 m / s ) , strain energy release rate was determined according.to Williams (l). A free vibrating torsional pendulum was employed for dynamic mechanical measurements.

Results and Discussion For all four types of the composites the yield stress a decreases linearly with the filler volume fraction v^ (Fig. 1) ; the slope depends on the matrix polymer and on the cooling rate.

Fig. l. Effect of the filler fraction Vf on the tensile yield strength Oy . 1 • • PE-C-F ; 2 o o PE-C-S ; 3 • • PP-C-F ; 4 • • PP-C-S .

547

log G' Pa 9 PE/CaC0 3 8

XT' • logG' // / Pa 7

W .

x

W\ V ^ PE

N i 0

-100

°c

100

Fig. 2. Temperature dependence of the storage G' and loss G'' moduli of PE and of PE composite with 20 v% of CaC0 3 .

Fig. 3. Temperature dependence of the storage G' and loss G'' moduli of PP and of PP composite with 28 v% of CaCC>3 .

As theoretical dependences

Oy

vs.

Vf , calculated for various

model space distributions of filler particles (2), differ from our experimental results, we can infer that relaxation of local stresses and plastic flow of the matrix adjacent to filler particles take place with increasing tensile strain. The effect of the cooling rate and, consequently, of supermolecular structure is significant for PE-C, whereas for PP-C it is rather small (Fig. 1). Likewise, much greater differences and shifts of the temperature dependences of the storage and loss moduli (at temperatures below 2CTC) can be observed between samples F and S for PE-C rather than

548 for PP-C (Figs. 2 and 3). Comparison of Figs. 1 , 2 and 3 reveals that the decrease in

Oy

with

Vf

is the faster, the lower the

matrix compliance. A steeper decrease in the pendence at

V£im in diameter at a particle concentration in the composite of 5 $ w/w.Glass beads 2-20 }im in diameter, nonsurface treated, were supplied by Jablonec Glassworks, Czechoslovakia. Si lanized kaolin was in the form of platelets (2-15 Mm). The concentration of glass beads and kaolin was in the range 30-50 % w/w. Polyethylene and the model spherical fillers (silica and glass beads) were mixed on a double calender at 150°C for 15 rain.The kaolin containing composites were compounded in a WemerPfleiderer ZDSK 53 mixer at about 200°C and 200 rpm.

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

562

Fig.1. Polyethylene DMDJ 3^72 filled w i t h 5 a. slightly etched - 1000 x;

b. deeply etched - 3000 x.

of

silica

563

b . 3000 X.

564

F i g . 3 . P o l y e t h y l e n e DGH 5320 f i l l e d w i t h 50 % w/w o f ( s l i g h t l y etched) a . 1000 x ;

b. 3000 x.

kaolin

565

Fig.k. Polyethylene DGH 5320 containing 50 fi w/w of kaolin (deeply etched) a. 1000 x;

b. 3000 x.

566 Compression moulded samples were polished (hand sanded and hand polished using aluminium oxide suspension) and then etched in xylene vapours at 100-110°C for 3-15 min.Micrographs were obtained on a Jeol JSM U3 scanning electron microscope.

Results and Discussion Fig.1 shows the system PE-silica.The contours of individual spherulites are clearly seen when the sample is only slightly etched (Fig.la).A deeply etched sample (Fig.lb) shows the internal structure of the spherulites.The size of the spherulites lies within 10 to 20 pm (cf.ref.2,3).Polyethylene filled with glass beads is shown in Fig.2.The matrix exhibits a lamellar grainy structure, the grain size varying from 2 to 8 yaa. Glass beads covered with polymer appear as clear regions. Kaolin filled polyethylene (Fig.3a,b) shows a grainy structure similar to that of the polyethylene filled with glass beads.Veil developed lamellar regions resembling the texture of coarse fabric (cf.ref.3) are seen in the case of samples etched at higher temperature (Fig. *ta,b). For the examined composite systems, the structural model by Chacko et al.(3) may be accepted;i.e.the filler particles are surrounded by a nonuniform interfacial layer and the matrix itself consists of subspherulitic structures.The interface in our composites is created from polymer links (l,cf.ref.3»^ "rodlike" and "random" morphology).The polyethylene matrix consists of subspherulitic lamellar structures when glass beads or kaolin are used as filler. Only the system silica-polyethylene contains spherulites,the interfiller particle distances being sufficient to permit their development.

Conclusion Polyethylene filled with kaolin and model composite systems (PE-silica, glass beads) have been examined.lt has been found that polyethylene has a spherulitic morphology when filled with

567

5 $> w/w of silica (particle size about 0.3 >im).Vhen filled with glass beads or kaolin (30-50 % w/w,particle size ranging from unity to tens of jam), no long-range ordering on the scale of spherulites has been observed.Instead,a pronounced subspherulitic lamellar morphology was noted,particulary in the system polyethylene- kaolin.The spherulites size was in the range 10-20 jim, while that of the subspherulitic structures was only several yun.

Acknowledgement The author wants to thank to PMM editor Prof,B.SedlaSek and Prof. P.H,Geil(University of Illinois)for their valuable advice.

References 1. Dolâkovâ,V. and F.Hudeöek.1978.Structure of filled linear polyethylene. J.Macromol.Sei. Phys. B15, 337-3^6. 2. Cole,J.H. and L.E.St.Pierre. 1978..The role çf interfacial energy in the heterogeneous nucleation of polyether crystallization. J.Polym.Sei. Polym.Symp. 6 3 , 220-235. 3. Chacko,V.P., F.E.Karasz, R.J.Farris and E.L,Thomas. 1982. Morphology of CaCO„-filled polyethylenes, J.Polym.Sei,Polym. Phys. 20, 2177-2195k. Go,A., L.Mandelkern, R.Prudhomme, R.S.Stein. 1974. Light scattering studies of narrow fractions of PE. J.Polym.Sei. A-2, 12, 1U85-1U90.

DETECTION OF INTERFACIAL DEBONDING IN PARTICLE-REINFORCED COMPOSITES

E.A.A. van Hartingsveldt Akzo Corporate Research Department Arnhem, The Netherlands

Summary The mechanical properties of particle-reinforced composites are strongly dependent on the nature of the interphase between matrix polymer and filler. The volume-strain method presented here is able to detect the moment of debonding of the filler particles. Combination of this technique with hysteresis experiments has shown, that the elastic behaviour of a freshly prepared composite may not be affected even in case of very poor adhesion. However when debonding has taken place during a previous deformation, as detected by the volume-strain method, the poor interfacial adhesion does change the elastic constants of the composite significantly .

Introduction To a significant degree the mechanical properties of composite materials depend upon the adhesion between the filler and the matrix. Stresses acting upon the matrix are only transferred to the filler when the interfacial coupling is sufficient. Sometimes the coupling can be enhanced by the addition of a so-called adhesion promotor or coupling agent. Well-known is the use of silanes for the improvement of the adhesion of glass to different polymers (1 ) .

An essential tool for the development of new particle-reinforced plastics is a technique, that is able to give a valuation of the quality of the interfacial adhesion, thereby offering the possibility to check the effect of a coupling agent. The volume-

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

570 strain method, presented here, furnishes the opportunity for a semi-quantitative description of the adhesion and is relatively easy to perform.

The Volume-Strain Method All materials show an increase in volume when they are subjected to a tensile test. In the elastic region this property is expressed by the Poisson's ratio of the material. When the volume changes are very small, which is true for rubberlike materials,the Poisson's ratio

(v)

approached 0.5. For most polymers in the glassy

state v lies between 0.3 and 0.4. However, when a composite is strained and the reinforcing particles start to debond, voids will grow at the interface thereby creating an extra volume increase. This will be detected by the volume-strain method. This technique has been described using a dilatometer filled with gas (2) or liquid

(3). In this study the volume changes of the sample

are registered by the continuous measurement of the three dimensions of a test bar during unidirectional straining in a tensile tester. A Zwick analog extensometer is used to measure the elongational strain, while the contraction of the width and the thickness of the bar is monitored using two transverse strain sensors (Instron-2640) , which makes a very accurate determination of the volume-strain possible.

Materials The matrix polymer chosen for this study is polyamide-6 (Akzo (!) Plastics, Akulon K123). The filler particles are glass spheres with a diameter smaller than 50 ym (Potters Ballotini 3000, no surface treatment). The coupling agent used is yaminopropy1triethoxysilane. The samples were compounded on a 19 mm Gottfert single screw extruder and injection moulded after standard drying procedures.

571

Results and Conclusions Polyamide-6 filled with 30 wt% glass spheres has been analysed using the volume-strain method. Moreover,the influence of the addition of 0.3% aminosilane has been determined. Stress-strain as well as volume-strain curves are presented in Figure 1 (samples conditioned: 23°C/50% R.H.). Up to 1.5% strain the samples behave in the same way, but at that point the stress build up in the sample without silane nearly stops and at the same strain-level the volume-strain of that sample starts to increase far more rapidly as a result of the creation of interfacial voids. It is striking, that there is no effect of the coupling agent at small strains. Elastic constants like Young's (or E-)modulus and Poisson's ratio do not change. The elastic behaviour of both materials can be described very well using the well-known equations derived by Kerner (4) and Chow (5). However, this is only valid for a freshly prepared sample that has not been deformed during a previous experiment. When the composite with the poor adhering glass spheres is deformed to a strain-level exceeding 1.5%, both the E-modulus and the Poisson's ratio decrease dramatically.

STBISS (KPi)

40

VOL0ME

•—•SIMPLE COlTiimSG 0.31 C.A. •—• SAMPLE fITHOUt C.i.

STRAIN

30 -

-0 75

M -0 50

-0 25

0

Z

3 5 ELOHGATIONiL SIHAI5 (!)

6

Fig. 1 . Stress-strain ( ) and volume-strain ( of polyamide-6+30% glass spheres.

7

) behaviour

572

PRKSTSAIB (!) Fig. 2. Polyamide-6 + 30% glass spheres: Dependence of the Young's modulus on the prestrain (i.e. the maximum strain during a previous experiment).

PHsmiK (J) Fig. 3. Polyamide-6 + 30% glass spheres: Dependence of the Poisson's ratio on the prestrain (i.e. the maximum strain during a previous experiment).

573 This is demonstrated by a number of hysteresis experiments. A test bar of the composite is deformed to a fixed strain-level and after removal of the stress the experiment is repeated, whereby each time the extension is increased. Stress-strain as well as volumestrain curves are recorded continuously. In this way the dependence of the E-modulus and the Poisson's ratio on the prestrain

(i.e.

the maximum strain the material has had during a previous experiment) can be determined. The results are presented in Figures 2 and 3 (samples as moulded). The following conclusions can be drawn: a) Young's modulus and Poisson's ratio

of unfilled polyamide-6

are hardly affected by a previous deformation. b) It is obvious, that in the composite without silane debonding of glass spheres takes place mainly between 2% and 6% strain. The E-modulus is even lower than the E-modulus of unfilled polyamide-6 when the composite has been strained to more than 4% in a previous test. c) The sharp decrease of both elastic constants is almost absent when a silane coupling agent is added. Figure 4 shows two subsequent volume-strain curves of the polyamide/glass composite without silane.

VOLUHE STRilll (*)

1.5

1.25

1 0.75

0.5

0.25

00

1

2

3

4

5

6

7

6

ILOIGitiOHU STBAIN (J) Fig. 4. Polyamide-6 + 30% glass spheres: Two subsequent volume-strain curves: A)new sample strained to 7%, B) second test after removal of the stress.

574 First a new sample is deformed up to 7% strain

(curve A). After-

wards the stress is removed and 1.8% permanent strain remains, that causes no volume-strain. Apparently this permanent strain is the result of shear deformation. All the voids created during the test have no volume anymore. Still the interfacial adhesion has not been restored as can be concluded from curve B as well as from Figures 2 and 3. The volume-strain method is able to detect the moment of debonding of the filler particles. The combination of this technique with hysteresis experiments as described above furnishes the opportunity to estimate the strain interval in which debonding takes place and moreover to get quantitative information about the degree of debonding.

Acknowledgements The author wishes to thank Dr. J.J. van Aartsen, Dr. M.G. Northolt and Dr. D.W. Koetsier for the many stimulating discussions and their continued interest in this work.

References 1. Plueddemann, E.P. 1982. Silane Coupling Agents, Plenum Press, New York. 2. Farris, R.J. 1964. J.Appl.Polym.Sci. 8, 25. 3. Coumans, 11.J., D. Heikens. 1980. Polymer 21, 957. 4. Kerner, E.H. 1956. Proc.Phys.Soc. B69, 808. 5. Chow, T.S. 1973. J. Polym. Sci. , Phys.Ed. 16, 959.

THREE-FIBRE METHOD FOR MEASURING GLASS FIBRE TO THERMOPLASTIC BOND STRENGTH

P.A. Jarvela, P. Tormala Tampere University of Technology, Institute of Plastics Technology, P.O. Box 527, SF-33101 Tampere, Finland P.K. Jarvela Oy Partek Ab, Development Centre, SF-21600 Parainen, Finland

Introduction The properties of fibre reinforced polymer composites are mainly controlled by the strength of the bond between the matrix and the reinforcement. It is at this interface where stress concentrates, no matter whether it arises from differences in thermal expansion coefficients, from loads applied to the structure, or from cure shrinkage and crystallization. The interface may also serve as a nucleation site, a preferential adsorption site, and a locus of chemical reaction. Thus it is no surprise that considerable efforts have been made over the last 20 years to gain better understanding of the interface and to device ways for controlling or even modifying it (1 ). At present we know at least two effective ways to increase bond strength at the interface: surface treatment of the fibres and chemical modification of the polymer matrix (2). Results from adhesion tests between matrices and fibres have recently become available. These tests can be divided into two groups: massive-scale tests performed on e.g. boards, cylinders and rods (3, 4), and micro-scale tests on single fibres (2, 5-6). Our aim in this study was to device a test method for determining bond strength between a single fibre and thermoplastic matrix.

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

576 Fibre Pullout Test For short fibre reinforced polymer composites the properties of the fibre-polymer interface and the aspect ratio of the fibre i.e. the ratio of length to diameter - are important factors that control the mechanical properties of the composite (1, 2). We know that if the length of a fibre with a given diameter falls short of a critical value, the transfer of stress from the matrix to the fibre will be insufficient. The simplest model for estimating critical fibre length 1 lc=6

f d

/2T

where d is fibre diameter,

is (2)

B

(1)

6 f is the ultimate fibre strength,

and Tg is the interfacial bond strength. Both the interfacial shear strength x B and the critical length 1

can be esti-

mated by eq. (1) and the single fibre pullout test, in which different lengths of the fibre are embedded in the matrix and the force required to pull the fibre out is then measured.

Experimental Materials The study concentrated on type A glass fibres. The fibres were manufactured by drawing from liquid glass without surface treatment. All tests were conducted with horizontal fibres of 180 ym in diameter. The diameter of the vertical fibres was 90 ym. The bonding material consisted of two thermoplastics: a high density polyethylene (Hostalen GA 7260, supplied by Hoechst) and polypropylene (Novolen 1100 TX, supplied by Basf).

Test equipment and sample preparation The method, called the three-fibre method, is shown schematically in Fig. 1. The vertical fibre is pulled out from a drop of adhe-

577 sive between it and the horizontal fibres whilst the movement of the horizontal fibres in the direction of the pull is prevented. This principle is very simple and can be applied to different fibre/adhesive bonds.

Fig. 1. Schematic representation of the three-fibre method used to determine bond strength between fibre and thermoplast. Fig. 2 shows a cross-section of the testing equipment used. The fibres were pulled out from the polymer matrix by a J.J.Lloyd T5003 tensile testing machine. The experiments were performed at 23 °C. The speed of the moving clamps of the machine was 1 mm/min. The samples were prepared from fibres which had been stored in a dark, dust free place. The thermoplastic was used in the form of a 25 /jm thick film. The vertical fibre was threaded through a very small hole in a small piece (0 1 mm) of the thermoplastic

578 film. The distance between the horizontal fibres was then reduced, thereby bringing the horizontal fibres into contact with the thermoplastic film, after which we checked that the fibres did not touch each other. The pullout test assembly was then placed in a preheated oven (approx. 250

C) for 12 minutes. Part of the

specimens were prepared in an air atmosphere and part in a nitrogen gas atmosphere.

Fig. 2. A cross-section of the equipment used for preparation and testing of the three-fibre bond

Results and Comments The three-fibre method measures the pull resisting strength of the bond when the vertical fibre is pulled at a constant speed. In practice a breaking in the system measured can occur either as a fracture of the vertical fibre or as a breaking of the bond. In this test a fracture of the vertical fibre is considered undesirable, because it indicates that the bond has greater tensile strenght than the vertical fibre.

One part of the measuring process is evaluation of the shear strength between matrix and fibre. Shear strength t b is calculated by eq. (2) t

B

= F

max

/irdD

(2) * '

where Fmax is the highest recorded value of the pull resisting force, d is the diameter of the vertical fibre and D is the dia meter of the horizontal fibre. We used fibres with a constant r diameter, which makes it possible to use F max instead of in comparing bond strengths measured from different T

matrix materials. Fig. 3 shows the maximum value of the recorded force (F ) max required to pull out fibres from HDPE and PP matrices.

1,6

1,2

0,8

0,4

0,0 I

1 A

1 B

1

1

C

D

Fig. 3. The force (F ) required to pull out fibres from HDPE and PP matrices™ 3 ^: HDPE (in air atmosphere, 250 C, 12 min ); B: HDPE (in nitrogen atmosphere, 250 C, 12 min ); C: PP (in air atgosphere, 250 C, 12 min); D: PP (in nitrogen atmosphere, 250 C, 12 min ).

580 The microscopic examination of the bond was performed using an ISI 40 scanning electron microscope (SEM). Figures 4 and 5 show bonds tested by the three-fibre method.

(b)

(c)

Fig. 4. Scanning electron microscope pictures of a glass fibre/HDPE bond: (a) the bond during pull (first fractures have appeared in the matrix); (b) a vertical fibre after the bond has completely fractured; (c) horizontal fibres after complete fracture.

581 The samples were prepared either in air atmosphere or in nitrogen gas atmosphere. In the case of samples prepared in nitrogen gas atmosphere stronger forces are required to pull out fibres from the matrix. The oxygen of an air atmosphere reduces bond strength at increased temperatures; therefore, it is preferable to prepare samples in an inert gas atmosphere. The melting index (MI) of thermoplasts should also be taken into account; the higher the MI, the shorter melting time and lower temperature is required, and vice versa. The three-fibre method has earlier been applied succesfully for measuring the bond strength of a fibre-resin bond (6). Now we applied the same method for a fibre-thermoplast bond. Both previous preliminary tests and our tests go to show that this method is a very fast and relatively accurate indicator of the strength of fibre-thermoplast bonds.

582 References 1. Kardos, J.L. 1984. Composite Interfaces: Myths, Mechanisms, and Modifications. Chemtech (July 1984), 430-434. 2. Westerlind; B., M. Rigdahl, H. Hollmark and A. De Ruvo. 1984. Interfacial Properties of Regenerated Cellulose Fiber and Thermoplastic Systems. J.Appl. Polym. Sci. 29, 175-185. 3. Yamaguchi, Y. and Amano, S. 1974. Effect of Coupling Agent on the Adhesive Strength between Glass Plate and Resin and on the Tensile Strength of F.R.P. SPE, Ann Tech Conf, San Francisco, CA, USA, May 1974 Paper 32, 452-455. 4. Laws, V. 1982. Micromechanical Aspects of the Fibre-Cement Bond Composites 21» 145-151. 5. Favre, J.P. and Merienne, M.-C. 1981. Characterization of Fibre/ Resin Bonding in Composites Using a Pull-out Test. Int. J. Adhesion and adhesives 311-316. 6. Jarvela, P., K.W. Laitinen, J. Purola and P. Tormala. 1983. The Three-Fibre Method for Measuring Glass Fibre to Resin Bond Strength. Int. J. Adhesion and Adhesives 2 141-147. 7. Bentur, A., S. Mindess and S. Diamond. 1985. Pull-out Processes in Steel Fibre Reinforced Cement. Int. J. Cement Composites and Lightweight Concrete 1_, 29-37. 8. Chua, P.S. and M.R. Piggott. 1985. The Glass Fibre-Polymer Interface: I - Theoretical Consideration for Single Fibre Pullout Tests. Composites Science and Technology 22^, 33-42. 9. Chua, P.S. and M.R. Piggott. 1985. The Glass Fibre-Polymer Interface: II - Work of Fracture and Shear Stresses. Composites Science and Technology 22[, 107-119.

INFLUENCE OP THE VISCOELASTIC PROPERTIES ON THE BONDING STRENGTH OP METAL-POLYMER COMPOSITES

A. Bauer, C. Biechof Academy of Sciences of the G.D.R., Institute of Polymer Chemistry "Erich Correns", 1530 Teltow, Kantstrasse 55, G.D.R.

Introduction The coating of metals, especially steel strip and aluminum, with polymers has gained an ever increasing extension in the past few years. This development can mainly be attributed to the properties of polymers affording protection against corrosion by attack of environmental media. The performance and service-life of metal-polymer composites depend largely on strong and permanent bonding of the components in the interface layer. For the determination of the adhesion strength between metal and polymer many methods are used, for instance cross hatch test, peel test, pull-off test, Erichsen cupping test and some others. With respect to these various possibilities of testing the following question arises: do the different methods result in consistent statements or do the viscoelastic properties of polymers at different loading stresses exhibit considerable differences in the numerical values?

Materials In our experiments PVC-coated steel strips were used. As PVC does not adhere to metals, it is necessary to use a primer. Therefore, the composite is composed of three layers with two interface layers between them (Pig. 1). The formation of good adhesion in these two interface layers is very important. In our tests we used different PVC-plastisols like Plastisol 1 (50 % PVC + 50 % DOP), Plastisol 2 (60 % PVC + 40 % DOP) and Plastisol 3 (70 % PVC + 30 % DOP) and also two primers of different para-

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

584

Polymer protective layer 50- 200jjm

- - - y

/interface layer between polymers

— Primer • Interface layer between primer and metal "Metal Pig. 1. Scheme of the PVC-metal composite meters. Films cast from the different PVC-plastisols were tested for their stress-strain behaviour (Table 1). The stress-strain-behaviour is shown to depend on the proportion of P7C and DOP, Young' modulus being the most responsive to the PVC/DOP ratio. Table 1.

Stress-Strain Values of Different PVC-films (0.3 - 0.6 mm thick)

PVC/DOP

50/50

60/40

Break stress, MPa

13.6 + 0.6

16.1 + 0.7

19.2 + 0.9

Break strain, %

350+11

250

250

Young's modulus, MPa

8

Modulus at break, MPa

3.2

+ 0.2

34 5.2

+ 10 + 0.6

70/30

87

+ 24 + 0.9

6.4

Testing methods For our tests the following three methods were chosen. In the pull-off test (1) the samples were centrically stuck between two stamps of metal and then they were stressed without bending moment using an Instron machine. The bonding strength was measured and the break type was estimated. Because under conditions of practical application the coated strips are strained, we have strained our specimens for 18 % before testing by the pull-off test. Using the Erichsen cupping test, a ball with a diameter of 20 mm

585 was pressed into the bottom side of the coated strip with a defined depth. To account for the elasticity of the PVC-coating, it is necessary to cut two crosses in a distance of 5 nun into the PVC before testing. After the cupping test, the coatings were separated from the strips using break by hand. The Erichsen value is defined as the length of the uncovered strip and it is a measure of the quality of adhesion. In Fig. 2, correlations between the Erichsen value, the cupping depth and the composition of different samples are demonstrated. The Erichsen value de-

Depth of the c u p x

4 mm

o

6mm

•

8mm

10

i 4

5

Sample number

Pig. 2.

The Erichsen cupping test.

pends directly on the depth, while the difference between the Erichsen values of different samples is independent of the depth. By a graphical-analytical method the depth was found to be 7.5 am, being equal to the strain of the sample for 18 %. In the peel test (2, 3) the force required for peeling the coating off the strip is determined. At a peel angle of 180° the peel force is calculated by the formula advanced by Kaebble (4) P = hb a 2 /4E where h is the thickness, b is the width, a is a not defined

586

stress of the coating and & is the Young's modulus of the coating films.

Results The bonding strengths obtained by the pull-off test are summarized in Table 2. Two kinds of samples were used, namely coated strips and PVC-films stuck between the two testing stamps. The poor primer (SZ B) caused adhesion breaks and low bonding strengths. By applying a good primer (SZ C) the break occurred in the PVC-layer and the strength values were in agreement with the film strengths. The data gained by the pull-off test are affected by the tensile properties of the coating. In the Erichsen cupping test tensile strength and bonding strength are superimposed and, thus, at the crosscut of the coating a bending moment is caused with the factor EI = Ebh 3 /12 Table 2.

Bonding Strengths, Erichsen Values and Feel Strengths of Different Coated Strips

PVC/DOP Bonding strength, MPa

Erichsen value, mm Peel strength, MPa Peel distance, mm

50/50

60/40

70/30

SZ B

4.6 ± 0.8

7.4 + 0.2

SZ C Film

6.7 + 0.8 6.4 + 1.2

7.9 ± 0.9 9.3 + 0.7 8.2 + 3.0 10.6 + 2.1

SZ B

8.1 + 1.2

SZ C

5.1 + 1.6

9.2 + 2.1 13.5 + 3.9 5.7 + 1.9 8.2 + 1.8

4.9 + 0.3

6.5 + 0.8

8.0 + 0.9

77 + 15

66 + 6

95 + 5

6.7 + 1.4

For Plastisol 3 and Plastisol 1 films (having, respectively 0.335 and 0.633 mm in thickness) the factors EI are in agreement (0.775 and 0.666); the peel forces (1.91 and 2.12 kp) and the Erichsen values (7.9 + 1.4 mm and 8.1 + 1 . 2 mm) are also in conformity. In the Erichsen cupping test the elastic modulus and the transferable peel forces (breaking Btress and cross-section)

587

are decisive. It is necessary to use only equal coatings for the testing of primers. Relating the peel forces and the peel distances found with different coated strips, plausible quantities were obtained by using the Kaebble formula only, provided that for such calculation the modulus E and the stress a at the breaking point are used. These results show that, if the adhesion strength is nearly equal to the cohesion strength, the peel force depends on the breaking strength of the coating and on the modulus near to the breaking point. In this case information about the adhesion is not provided by the peel force, but only by the peel distance. The following Fig. 3 shows an attempt to elucidate the main parameters influencing the different test methods.

Testing nethod

Stresses

T

Pull-off teat

.a.

TT

Peel test

Brichsen cupping test

Pig* 3.

r

Range of the stress-strain curve account for the numerical value

tensile forces on the coating layer with 'obstruction or the cross-construction

- breaking stress

tensile forces at the peeled coating, bending aoaent and bending forces at the peeling place

- aodulue above yield point - breaking stress

tensile and bending forces during the cupping produce a break off at the cutting placet at the end of the test a peeling follows

- Young's modulus - modulus above yield point - breaking stress

Stress occurring in the coating in dependence on testing methods.

588 MPo

~5

10

mm

15

E n c h s « n value

Pig. 4.

Inadequate correlation of Erichaen data obtained for unequal PVC-coatings.

Conclusion Good adhesion is characterized by a high value of pull-off test and a low value of Erichsen cupping test. If different coatings are used for testing primers, only test data obtained with equal PVC-coatings should be correlated otherwise wrong results could be obtained (Pig. 4). The two full lines SZ B and SZ C show correlation between bonding strength and Erichsen values for different coating films, but these correlations are not accurate. Only the dashed lines reflect the true relations between the values obtained by the two methods.

References 1. A . Bauer, C. Bischof. 1983. Plaste u. Kautschuk ¿0 4, 208-211. 2. J. I. Gordon. 1963. J. Appi. Polymer Sci. 2» 643-665. 3. A. N. Gent, G. R. Hamed. 1977. J. Appi. Polymer Sci. 21., 2817-2838. 4. D. Kaebble. 1971. Physical Chemistry of Adhesion. Wiley-Interscience, New York-London-Sydney.

CAFOD - COMPUTER-AIDED FIBER ORIENTATION DETERMINATION IN COMPOSITES

V. Djakovich, S. Fakirov and L. Christov University of Sofia, Laboratory on Structure and Properties of Polymers, 1126 Sofia, Bulgaria

Introduction The mechanical properties of the discontinuous ("short") fiber reinforced thermoplastics depend on many factors, including the fiber orientation of the short fibers and its distribution

(1-4). The

orientation of fibers plays a particular role in cases when it is not uniform with respect to the mold fill direction (MFD), as found for injected molden plaques of poly(ethylene terephthalate) (PET) (5) and nylons (6). The laminate microstructure consist of three observable layers: two surface layers with fibers highly alignedin the MFD, and a core with fibers lying mainly transversely to both the MFD and the plaque thickness direction (6,7). The knowledge of the fiber orientation and other chemical and structural characteristics, makes possible to predict quantitatively the mechanical properties of the polymer composites

(1,2).

A simple and new technique for such a direct determination was recently proposed in ref.(8) based on the fact that the shape of the fiber cross section on the plane perpendicular to the desired direction (e.g. orientation direction) depends on the orientation angle. For an eliptical shape with minor (r) and major (R) radii one obtains cos $ = r/R (8). Using micrographs of the polished sample cross sections of commercial PET reinforced with 45% short E-glass fibers and assuming a planar fiber distribution, the following parameters have been determined

(8): the - values, the o-

rientation parameter f , the angular distribution

(histograms),

the number and the sizes of the layers with more uniform fiber orientation, and the transition zones between them. In contrast to the previous techniques (7,9,10), this method makes possible evaluating of the above mentioned properties in a desired section or in the whole volume (not only in one plane or in a thin slide) of the sample studied.

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

590 The purpose of the present work is to extend the described technique (8) by introduction of a new parameter, sin tp, derived from the slope of R vs. x dependence. This parameter makes possible to evaluate the deviation of the glass fibers from planes parallel to largest surfaces of the sample. A computer program is elaborated which allows one to obtain a rapid quantitative characterization of the fiber orientation in fiber reinforced composites.

Experimental The samples used are commercial short E-glass fiber reinforced PET of type Rynite (E.I. du Pont de Nemours and Company, Inc.). The weight fraction of the fibers was 45 percent and their average diameter 10 ym with a length of about 200 ym (7). Details concerning the molding conditions are given elsewhere (5). All the measurements in the present work were performed on micrographs used for the same purpose in the previous work (8). They are taken from the entire cross section of the injection molded, rectangular plaques (127 x 7.62 x 3.18 mm) (7). The micrographs were step-wise magnified (ca.10x) in order to obtain fiber cross sections of at least 8 to 10 mm in diameter. From these micrographs, the following variables were measured: r and R, the coordinates x and y of each fiber cross-section center, and the slope of R with respect to x. The slope of R allows one to calculate the orientation parameter of fibers, which do not lie in the plane assumed for the other fibers. From the orientation angle values the orientation parameter, f p , proposed by Hermans (11) for describing the orientation in crystalline polymers was calculated. For the planar distribution being considered here, the "Hermans' orientation parameter", f p , is of the form: 2

f. P

2 - 1

where 2

= ¡^(((K) cos^./jNl^.)

(1)

591 and represents the angle between the individual fibers and the primary axes (= MFD) , and N(

47

sin 0

0 .500

0 .845

0 .325 0 .0247 -1 .542

0 .243

0.103 0.224

-1 .318 0 .442

2.064 3.085

C. As seen from Table 1 the short glass fiber orientation in molded PET plaque is characterized by: 1. Inhomogeneous fiber distribution with respect to sample thickness . 2. A well expressed sandwich structure can be seen in accordance with previous observations (6-8) . 3. The dominating angle values of and S are very close to 0", followed by 50° (for (9

i n < N i n

m

k O i o i N n c n C O C T— *— c o T C M

T—

•

o

C M i n •

o

I

T—

o

«"

O

1

o

O

t o r— O

o

m

T—

C M n

T

o r» i n k

00

o CTI

ÍH increases

from 20" up to 62'). 3. The central core of the sample, occupying 1/5 of the total sample thickness, demonstrates the average orientation angle between 62 and 64". 4. Again three principal zones are observed: surfacial, transitional, and central ones, in accordance with recent investigation (8) .

The results obtained demonstrate the lamellar character of the composite material studied; this must be taken into account when considering the mechanical behaviour of such type of composites.

References 1. McCullough, R.L. 1979. Influence of Microstructure on the Thermoelastic and Transport Properties of Particulate and Short-Fiber Composites. In: Mechanics of Composite Materials, John Wiley & Sons, New York. 2. Wu, C.T.D., R.L. McCullough. 1977. In: Development in Composite Materials-I (G.S. Holistr, ed.). Applied Science Publishers, JTD, London, Ch. 7. 3. Whitney, J.M., I.M. Daniel, R. Byron Pipes. 1982. Experimental Mechanics of Fiber Reinforced Composite Materials. The Society for Experimental Stress Analysis, Brookfield Center, Connecticut. 4. Pipes, R.B., R.L. McCullough, D.G. Taggart. 1982. Polym. Compos. 3, 34. 5. Wetherhold, R.C., W.A. Dick, R.B. Pipes. 1980. Thickness Effect on Material Properties in a Glass/Thermoplastic PET Injection Molding Compound. SAE Technical Paper 800812. 6. Obiego, G., D. Yilmaz. 1983. Kunststoffe 73, 83.

595 7. Friedrich, K. 1982. Microstructure and Fracture of Fiber Reinforced Thermoplastic Polyethylene Terephthalate. Fortsch.Ber. VDI-Z Reihe 18, Nr 12. VDI-Verlag, Dusseldorf. 8. Fakirov, S., C. Fakirov. 1985. Polym. Compos. 6, 41. 9. McGee, S.H., R.L. McCullough. 1982. An Optical Technique for Measuring Fiber Orientation in Short Fiber Composites. Proc. U.S.A.-Italy Joint Symp. on Composite Materials, 1981, Capri, Italy. Plenum Publ. Corp., New York. 10. McGee, S.H., R.L. McCullough: J.Appl.Phys. (in press). 11. Hermanns, P.H. 1946. Contributions to the Physics of Cellulose Fiber. Elsevier, Amsterdam.

HOMOGENEITY OF POLYMER COMPOSITES

Frantisek Rybnikar Technical University Brno, Faculty of Technology 762 72 Gottwaldov, Czechoslovakia

Introduction In a previous article (1) on the homogeneity of crystalline polymer samples we have discussed factors which essentially hinder one from preparing isotropic samples for structure and property investigations. It is the aim of this report to show that the situation is even worse for filled samples. Here, it is important to use not only an isotropic polymer but also a homogeneous dispersion of the filler in the polymer matrix. For polymer filling mineral fillers are used very often. The achievement of a good dispersion of solid fillers in the polymer depends on several factors, such as shape of filler particles (e.g. particulate in calcite, kaolin and chalk, lamellar in talc and mica, fibrillar in glass fibres or asbestos), size of filler particles and their tendency to aggregate or interact with the polymer. Usually, a completely random dispersion of filler particles in a polymer is desirable. The degree of filler dispersion depends on the mixing procedure, on the size and distribution of filler particles and on the wetting of filler by the polymer. It is evident that local differences in filler concentration or filler size distribution will bring about local differences in physical characteristics. In this respect there are especially critical and unwanted coarse filler aggregates, because they represent serious structural defects. The optimal size of particulate fillers is in the range of 1 to 100 pm which influences the choice of experimental methods for testing the homogeneity of filled polymers. In this article, we describe the application of two simple methods for testing the homogeneity of calcite particles dispersion in an isotactic polypropylene.

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

598 Experimental Materials. Commercial samples of isotactic polypropylene (PP) Mosten 58412 (Chemopetrol, CSSR) and its composites with CaC0 3 Durcal 2 (Omia, France) of the calcite type were used. Particle size of the filler was mainly in the range of 1 to 10 ym. The sheets were melt pressed from the granulate (220°C, 18 MPa). Contact radiography. For testing the homogeneity of mineral filler distribution in PP, a modified contact radiography method was used. In this procedure, 0.1 mm thick sheet of the composite was put on the X-ray film in a light tight envelope and the assembly was irradiated by X-rays (CuKa). After usual photographic procedure a contact radiogram is obtained showing the distribution of regions within the sample, which differ in absorption of X-rays. With our samples we were able to obtain information on the distribution of filler particles or their aggregates larger than 6 ym in diameter in an area of the sample about 4 cm in diameter. Smaller filler particles could not be distinguished from the background. In order to exclude eventual artifacts originating from the photographic procedure, we made always two independent radiograms from the same sample and took into account only features present in both radiograms. Magnification in the positive prints shown here is 9x. Local density distribution. If, in the given polymer composite, the density of the filler and of the polymer differs appreciably, one can use measurements of local density differences for testing the homogeneity of the filler particle distribution. In such a case, the particle size is not a limiting factor. The procedure used was as follows: The sample - a 0.1 mm thick sheet - was cut into a small parts of identical dimensions (2x1x0.1 mm) and the density of each part was determined by a flotation method. After this, it was possible to draw the density distribution curve, characterizing the composite homogeneity. A narrower density distribution means a better filler distribution, and vice versa. (Fig. 6 shows typical density distribution curves constructed from 3 measurements of the same samples. To get better resolution smaller samples should be used.)

599 Results and Discussion Fig. 1 shows contact radiograms of PP filled with continuous (a) and chopped

(b) glass fibers. The difference in filler dispersion

is apparent. Using thin samples we can get similar information also by light microscopy

(Fig. 1c). Unfortunately, the better

resolution of light microscopy is accompanied by several drawbacks: it cannot be used with opague or thicker samples and the investigated area is relatively small. Fig. 1 proves that with fillers of greatly anisometric shape (fibers, sheets), it is, a priori, impossible to obtain an isotropic composite. Contact radiograms in Fig. 2 illustrate differences in the homogeneity of PP composites with 40% (w) of talc. The original radiograms represent a relatively large sample area and, therefore, they may be used for evaluation of the filler content and its particle size distribution using methods of guantitative microscopy

(2). A serious limitation of the contact radiography is its

low resolution. Without special photographic materials we are not able to identify particles of a particulate filler smaller than about 5 ym. Nevertheless, also in the case of smaller filler particles the situation is not hopeless. As an example, the composite PP + 23% (w) CaCO^ was taken. Here, besides contact radiography and electron microscopy, also the indirect method, based on measurement of local density differences has proven its utility. Preliminary examinations have shown that: 1) there are many filler particles exceeding 10 ym (the upper limit of the particle size distribution) in the sample, representing filler aggregates, 2) there are local differences in distribution of the particle size and filler concentration, varying systematically with their location. Therefore, we have investigated how these differences are connected with the way of sample preparation. Fig. 3 shows typical parts of contact radiograms taken from the central and marginal part of a melt pressed sheet. Visual and quantitative evaluation of the particle size distribution

(Fig. 4) prove that a greater

number N of larger particles remains in the central region of the sheet. This is true not only for larger filler particles and aggregates; a similar trend was found also for small filler particles by evaluation of the electron micrographs of the same

600

Fig. 1. Contact radiograms of PP filled with continuous (a) and chopped (b) glass fibres, and (c) an optical micrograph of a sample similar to (a).

Fig. 2. Contact radiograms of two composites PP + 40% talc.

602

b

Fig. 3. Contact radiograms from the central (a) and marginal part of the composite PP + 23% calcite.

(b)

Fig. 4. Particle size distributions from the central and marginal part of the composite PP + 23% calcite.

603

Fig. 5. Electron micrographs of the composite PP + 23% calcite taken from the central (a) and marginal (b) part.

20

O CENTER • MARGIN

% 10

0 1.05 Fig. 6. Density distribution curves of the composite PP + calcite.

604

samples (Fig. 5). Density distribution curves of samples investigated by contact radiography and electron microscopy (Fig. 6) confirm the local anisotropy in filler concentration. In the central part the filler concentration is higher than in the marginal part of the sheet. It is clear that sample preparation by melt pressing results in a heterogeneous filler distribution. Flowing melt takes away smaller particles preferentially and the filler concentration decreases towards the marginal parts of the sheet. This local heterogeneity is not connected with differences in the filler concentration of individual granules, because similar results were obtained with sheets prepared from one granule only or from a mixture of a powdered polymer and filler. Sample preparation by injection moulding has a similar effect on the filler distribution in the polymer matrix. With growing distance from the gate the filler concentration decreases. This is illustrated for a composite PP + 20% (w) CaCC>3 sample from the spiral flow test, showing the decrease in local density as a function of the distance from the gate (Fig. 7).

1.060

o-^q.

o

o -o

1.056 •

g cm^-; 1.044 O.

1.040 -

o 0

10

cm

20

30

40

50

Fig. 7. Variation of density with distance from the gate in a sample prepared by injection molding (spiral flow test).

605 The above results show that there are simple methods for testing the homogeneity of filler distribution in polymer matrices. All these methods applied to a composite PP + calcite pointed out that, strictly speaking, preparation of an isotropic filled polymer sample cannot practically be achieved.

Conclusions The results described illustrate that there are available simple means for testing the homogeneity of mineral filler distribution in a polymer matrix. Contact radiography, electron microscopy and local density distribution are in agreement and show that in most cases the homogeneity of filler particle distribution in the polymer matrix is questionable. Moreover, the method of preparing composite samples from even homogeneous granulate may lead to local differences in filler concentration and particle size distribution .

Acknowledgement The author is indebted to M. Macikova and M. Stoklasek for their technical assistence.

References 1. Rybnikar, F. 1984. Plasty a kaucuk 21, 289. 2. DeHoff, R.T., F.N.Rhines. 1968. In: Quantitative microscopy (McGraw-Hill ed.). N.York.

MICROSCOPIC METHODS CHARACTERIZING THE DISPERSION IN MINERALFILLED THERMOPLASTICS

V.Svehlovd Chemopetrol, Research Institute of Macromolecular Chemistry, Tkalcovska 2, 656 k9 Brno, Czechoslovakia

Introduction and Experimental The optimal useful properties of composites cannot be achieved without a satisfactory dispersion of the filler in the matrix, i.e., the formation of undesirable agglomerates should be precluded.The following microscopic methods (cf.ref.1,2) have been considered useful for evaluation of the degree of dispersion in both polyethylene and polypropylene filled with glass beads, kaolin and calcium carbonate: transmitted-light microscopy (TLM) and scanning electron microscopy (SEM). , The composites to be examined were melt compounded using either a twin-screw extruder or'KO-kneader; i.e., the material was obtained in a pallatized form. Thin films (less than 50 pm) were prepared direct from the pellets, either by compression moulding using commercial equipment or in a laboratory by hot pressing bet2 ween a glass slide and a cover slip. Several cm of the film were first examined with the naked eye and the representative part was then taken for the TLM observation. To locate regions absent from large agglomerates a low magnification (lOOx) was first applied to examine the whole sample. The region without agglomerates was then inspected at greater magnification to observe individual filler particles for particle size distribution. The composites (especially those not containing agglomerates) may also be examined using SEM. The SEM may be applied either directly, to examine the surface of pressed specimens (no thin films being necessary) or after'pretreatment (3) of the sample plate by polishing and etching. The sample surface was initially hand sanded down to 600 grade emery paper then hand polished using aluminium oxide suspension and etched in xylene at 85°C for 15 min.

Polymer Composites © 1986 Walter de Gruyter & Co., Berlin • New York - Printed in Germany

608

Fig.l. TLM examination of a machine moulded film with k0% w/w of microground CaC0„ Durcal 2 (Omya Co. , France) ;good dispersion of filler without agglomerates - 100 x .

Fig.2, TLM examination of a sample identical with that in Fig.l except for degree of dispersion - 100 x a. original pellets: large and frequent agglomerates; b. after injection moulding: low number of small aggregates; c. after additional mixing in a Brabender plastograph: almost without agglomerates.

610

F i g . 3 - TLM micr a c o v e r s l i p co Yugoslavia a . 100 x ;

b.

6'tO x .

¡s s l i d e and 5S f r o m

611

Fig.k. SEM observation: polyethylene filled with 50 $ w/w of glass beads (polished surface) - 300 x.

Fig-5. SEM micrograph: polyethylene containing 50 % w/w of kaolin (polished and etched surface) - 300 x.

612

Fig.6, SEM examination: compression-moulded surface of the same composite as In Fig.1 - 300 x.

Results and Discussion Micrographs obtained by TLM are presented first. Fig.l shows a thin machine moulded sheet. A good dispersion of filler without agglomerates is seen in this case. Fig.2 represents a composite identical in composition with that in the previous case but with differing degrees of dispersion. The initial, very poor dispersion shown in Fig.2a was gradually improved (cf.Fig.2b and 2c) with increased mixing. At the same time, the toughness of composite enhanced significantly as shown in Table 1. Table 1. The Dependence of Composite Toughness on the Filler Dispersion Degree Sample

Impact strength kJm" 2

Fig. 2a Fig.2b Fis^c

50 110 180

Area fraction of agglomerates 1° 1.1 O.k 0.1

Heyn s diameter of agglomerates % 54 20 16

613

The dispersion degree is characterized by two parameters which were obtained as a result of stereologic analysis of photographs using a semiautomatic image analyzer Opton TGA 10: area fraction of agglomerates and average size of agglomerates expressed e.g. by Heyn's diameter (i.e. four divided by the ratio of agglomerate surface to its volume). Fig,3a,b shows a thin foil prepared by hot pressing between a glass slide and a cover slip at two different magnifications. The micrograph at lower magnification confirms a good filler dispersion without agglomerates.Filler particle size distribution is clearly visible at a higher magnification and may be evaluated using an image analyzer. This examination is in good agreement with that of the filler powder before compounding. The content of separate filler particles larger than 10 )ira influences composite toughness in the same way as particle agglomerates do. This finding is valuable because composite toughness is one of the most important useful properties. The following photographs were obtained employing SEM, Fig.k represents a polished plate. A polished and xylene etched sample is shown in Fig.5. The same composite shown in Fig.1 (TLM) is presented in Fig.6 examined by SEM but at higher magnification. Contrast between filler and matrix was enhanced by electric signal treat ment on a Jeol JSM TJ3 microscope. The last micrograph is not only the best one but the sample preparation does not require great effort. Because of these two reasons the last case is preferred to other ones. Both methods - TLM with thin films and SEM on compression-moulded surfaces - are fast and simple. Light microscopy was found to be the better and fully convenient method for the detection of particle agglomerates. Both methods are equally valuable for testing of particle size distribution, but in case the filler particles are less than about 1 jxm electron microscopy must be used.

Conclusion

The described microscopic methods are capable of detecting large

614

and small agglomerates of filler in thermoplastic composite materials containing different particle types (glass beads, kaolin, calcium carbonate), present at different concentrations from several to 70 ^ w/w. They are convenient for production testing because they are fast and simple. The methods may also be applied when information on filler particle size distribution is required. The qualitative visual evaluation of samples can be quantified using an image analyzer. Composite toughness strongly increases with decreasing content of agglomerates larger than 10 jim.

Acknowledgement The author wishes to thank to Prof.P.H.Geil (University of Illinois) and PMM editor Prof, B.Sedla&ek for their valuable advice.

References 1. Ess, J.W., P.R.Hornsby, S.Y.Lin and M.J.Bevis. 1984. Characterization of dispersion in mineral-filled thermoplastics compounds. Plast.Rubb.Proces.Appl. 4, 7-14. 2. Parfitt,G.D. 1978. Dispergierung von Pigmenten in Theorie und Praxis. Defazet 32, 322-331. 3. Svehlovd,V. 1984. Report VUMCH. Brno.

ABBREVIATIONS ABS BET CXA DRS DSC DTA ELS EPDM EPR ESR EVA GPC HEED IR LAM M w Mn MWS NMR PBPP PBFP PCL/SAN PE PET PMMA PP PVAC PVC PVF RIM SALS SANS SAXS SEM TEM Tg TSD TSL WAXS WLF

acrylonitrile-butadiene-styrene terpolymer Brunauer-Emmett-Teller method ethvlene-based polymer with vinyl acetate (DuPont Co.) diffuse reflectance spectroscopy differential scanning calorimetry differential thermal analysis equal load sharing rule ethylene-propylene-diene (monomer) electron paramagnetic resonance (= ESR) electron spin resonance (= EPR) ethylene-vinyl acetate copolymer gel permeation chromatography high-energy electron diffraction infrared spectroscopy longitudinal acoustic modes (Raman) molecular weight (weight average, number average) Maxwell-Wagner-Sillars effect nuclear magnetic resonance poly-bis- (phenoxy)phosphazene poly-bis- (trifluoroethoxy)phosphazene poly-e-caprolactone/poly(styrene-co-acrylonitrile) blend polyethylene polyethylene terephthalate polymethyl methacrylate polypropylene polyvinyl acetate polyvinyl chloride polyvinylidene fluoride reaction injection moulding small-angle light scattering small-angle neutron scattering' small-angle x-ray scattering scanning electron microscopy transmission electron microscopy glass transition temperature thermally stimulated depolarization thermally stimulated luminescence wide-angle x-ray scattering Williams-Landel-Ferry equation

AUTHOR

INDEX

Ana, P.

457

Kelen, T. 167 Kevin, G. 525 Khalturinskii, N.A. 311 Kiss, P. 473 Klason, C. 153 Kokta, B.V. 251 Kolarik, J. 283,537,545,553 Kowalewski, T. 141 Kozlov, G.V. 389 Koziowski, M. 297 Kryszweski, M. 141,531 Kubät, J. 153 Kuiera, J. 545 Kuniak, L. 353 Kuperman, A.M. 487,497

Bauer, A. 583 Bazhenov, S.L. 487 Beniska, J. 243 Berek, D. 323 311,487 Berlin, A.A. 457 Bertalan, G. 251 Beshay, A.D. 583 Bischof, C. 191 Boissard, R. de 465 Bolshakov, G.F. Bryk, M.T. 269 Burban, A.F. 269 Chartoff, R.P. 89,525 Chen, A. 347 Chow, T.S. 19 Christov, L. 589 Daneault, C. Djakovich, V.

Lednicky, F. 537,553 Lin, Y.G. 373 Lopour, P. 303

251 589

Enikolopyan, N.S. Eriksen, E.H. 89

Magonov, S.N. 199 Malczewski, J. 381 Marcincin, A. 243 Marosi, G. 457 Matena, V. 217 Maurer, F.H.J. 399,449 Michailov, M.C. 275 Milczarek, P. 531 Miller, J.D. 449 Minkova, L.I. 275 Miwa, M. 183 Molnar, I. 457 Morozova, T.M. 389

67

Fakirov, S. 589 Folkes, M.J. 33 Freitag, K.-H. 413 Fukuda, H. 51 Gähde, J. 431 Gaieski, A. 141 Geil, P.H. 347 Gorbatkina, Yu.A. 497 Grellmann, W. 233 Hardwick, S.T. 33 Hartingsveldt, E.A.A. van Hoffmann, H. 233 Horsky, J. 283,291 Houskovä, M. 207,217 Hrudka, J. 303 Hugo, J. 207,217 Ishida, H.

449

Jacobasch, H.-J. 413 Järvelä, P.A. 575 Järvelä, P.K. 575 Jelcic, Z. 363 Johnson, J.K. 525 Kamdem, P.D.

251

569

Natova, M.M. 389 Nemec, J. 479 Nezbedovä, E. 515 Novak, I. 3 23 Ohsawa, T. 183 Ovchinnikov, A.A.

67

Pääkkönen, E.J. 199 Pangrele, S. 347 Pascault, J.P. 373 Pegoraro, M. 105 Petruj, J. 123 Pigiowski, J. 297 Ponesicky, J. 515 Ponomarenko, A.T. 67 Popova, T.V. 311 Pritchard, G., 329 Puchkov, J.V. 487 Pukänszky, B. 167,553

618

Ranogajec, F. 363 Rätzsch, M. 413 Rusznák, I. 457 Rybnikár, F. 597 Sautereau, H. 373 Shevchenko, V.G. 67 Sidorenko, A.A. 465 §mejkal, F. 225 Sova, M. 507,515 Speake, S.D. 329 Spicer, H.G. 525 Stankevitch, R.P. 311 Sulc, J. 303 Svehlová, V. 561,607 Szijártó, K. 473 Takahama, T. 347 Takayanagi, M. 3 Tcheshkov, V.M. 389 Tchmutin, I.A. 67

Törmälä, P. 199,575 Trotignon, J.P. 191 Tüdös, F. 167 Vallois, A. de 191 Verdu, J. 191 Vesely, K. 123 Vilesova, M.S. 311 Vondrácek, P. 303 Wong, W.K. 33 Wozniak, L. 381 Wu, C.S. 347 Zachariev, G.Z. 389 Zahradnîckova, A. 123 Zelenski, E.S. 487 Zemanová, E. 243 Zemek, J. 323 Zilvar, V. 233 Zuchowska, D. 381

SUBJECT INDEX

ABS-copolymers (acrylonitrile-butadienestyrene) 276-280 - , deformability 277-280 - , relaxation transitions 277 - , thermomechanical analysis 276-280 Acetylation, epoxy resins 351 ff. Acoustic emission 223 Acrylamide - , adsorption of 390, 392-395 - , IR spectrum 392-394 Activators 292 Additives 457 Adhesion, adhesive 105,106,108,114,116,187,188 207, 432, 457,497, 499, 501-503, 505, 506, 512, 514, 569, 570, 574, 575, 583 - , interfacial 512, 514, 569,570 - , promoters 106-108,116,118, 432 - , strength 187,188, 497, 499, 501-503, 505, 583 - , tests 575 Adsorption 415 Agents - , blowing 225, 226, 230 - , coupling 374 Agglomerates 612, 613 Albumin, crosslinked 323-328 Alumina trihydrate - , treated with Ca-stearate 474, 476 - , treated with titanate 474, 476 yAminopropyl triethoxysilane 570 Amino resins - , crack resistance 389, 394 - , crosslinking of 389, 390 - , hydrolysis resistance 389-394 Anhydride cured epoxy resins 349 Anionic polymerization - , activators 292 - , initiators 232 - , poly cap rolactam 292 Aspect ratio 4, 75, 76,186-188, 576 Aramids 3

- , - , properties 325 - , - , role of spacer length 325, 327 Birefringence 40, 42 Block copolymerization/copolymers 8,14 Blowing agents 225, 226, 230, 231 Bonding strength 584, 586 Bonds, hydrogen 366 Bonds, van der Waals 218 Break strain 584 Break stress 584, 587 Brittle fracture 509, 510 Bridges of matrix 512 Broadening of transition region 102,103 Bubbles - , growth 225-288 - , shape 229 Bulk modulus 107 Butadiene-styrene elastomer - , relaxation transitions 275, 278 - , thermomechanical analysis 275-278

Calcium carbonate 141-148,167-175, 225-227, 232, 474, 475, 538-540, 545, 597 - , filler 207, 537, 553 - , powder 208, 218 - , treated with - , - , Ca-stearate 474,475 - , - , titanate 474,475 Calcium stearate 474-477 - , flow properties 474,477 - , impact resistance 475, 476 - , treated with - , - , alumina trihydrate 474, 476 - , - , calcium carbonate 474,475 Cap rolactam - , anionic polymerization 292 - , composites in situ 292 - , polymerized in rotating moulds 292 Carbon fibres - , carbon 37, 38,41 - , carbon/nylon 66 44 Azid functional silane 451 Catalyst cured epoxy resins 349 Cellular Beads - , composite 228-232 - , glass of283,373-375,378,379,399,403-407,410 - , core 227, 231, 232 Cellulose - , glass micronodules 363,364 - , fibre 156,162 Bending stress 419 - , flour composite 149 BET method 413 Cellulosic fillers 153 Binding of water at filler 413, 421 - , cellulose fibre 156,162 Biopolymer-SiCh composite 323-328 - , cellulose flour 156 - , human serum albumin-silicagel- , cellulose grafted 155,160 polyisothiocyanates 323, 324, 326, 327 - , - , preparation 323

620 -, -, -, -,

polyamide 6, of 159 polyethylene, of 156 polymethyl methacrylate, of 160 polypropylene, of 152,156 polystyrene, of 156 - , lignin 156 - , microcrystalline cellulose 156 - , microfibril 153,155,162 - , wood flour 156,158 Chain - , folding 615 - , length 349 ff - , - , of epoxy resins 349 - , mobility 364 - , - , of b u n d l e s probability model 52 Chalk as a filler 141-148 Charge carriers 365 Charpy impact strength 178,179,193, 233, 234, 549 - , impact resistance 194 - , impact tester 233, 234 Chemical blowing agents 225, 226, 230, 231 Coated steel strip 583 Cohesion - , defects 481 - , failure of fibre 497, 505 Color strength 243, 247, 248 Compatibility e n h a n c e m e n t of fillers 473 Compatibilizer 105,106 C o m p l e x b o u n d e d radicals 124,133 C o m p l e x fatigue process 480 C o m p l e x shear m o d u l u s 401, 407 Compliance - , effect of dispersed elastomer u p o n 222 Composites (see also Composites, individual polymers) 359 - , b i o p o l y m e r / S i 0 2 323-328 - , c a r b o n / n y l o n 66 44 - , cellular 228-232 - , elastomer modified 213 - , glass fibre 507, 512 - , graphite 292 - , hybrid 51-53, 59-61, 63, 64 - , in situ 291 - , m o l y b d e n u m disulfide 292 - , polyamide, of 291 - , polycaprolactam, of 291 - , polyester/glass laminate 335, 342 - , polypropylene, of 207, 217 C o m p o u n d i n g t e c h n i q u e 155,156 Concentration - , factor S C F 52 - , stress 52, 53, 57-59, 62-64 - , threshold 72 Conducting polymer composites 67 if - , electrothermal analogy 69 ff

- , generalized conductivity 69 ff - , organic material electronics 67 Conductivity 72 - , local 74 - , metallic 85 - , periodic fields, in 77 Conformation of macromolecules 43 Contact p h e n o m e n a 79 Contact radiography 599 Cooling rate 546 Copolymers - , ethylene-propylene 208 - , ethylene-vinyl acetate (EVA) 399, 4 0 3 - 4 0 8 - , methyl methacrylate-styrene t h e r m o m e c h a n i c a l analysis 280, 281 - , - , compatibility with PVC 280, 282 - , methyl methacrylate-butadiene-styrene 266-282 Copolymerization 8, 253 Coprecipitation 5,14 Core, cellular 227, 231, 232 C o r r e s p o n d e n c e principle 401 Cost, grafted wood fibre/PMMA composites 266 Coupling agents 159,193, 283, 374, 473, 4 7 7 , 4 7 8 - , silane 193 titanate 193 C o x - C h o w equation 19, 20 Crack 7, 207, 332,434, 480 - , blunting 208 - , initiation in polyester resins 332 - , osmotic 332 - , pinning 207 - , propagation 483 - , resistance 389, 394 Craze 14,454 Critical - , fibre length 44, 45, 47,183,184,187-190, 576 - , - , interfacial microstructure, effect on 44 - , - , m e a s u r e m e n t technique 45 - , - , polypropylene, of 81,184,187-191 - , micellar concentration 418, 423 - , mixing time 413-415, 423-428 - , point of liquid crystal formation 5 Cross hatched morphology 583 Crosslinking - , crosslinked a l b u m i n s 323-327 - , - , comparison with biopolymer/SiC>2 composite 325-327 - , - , preparation 324, 325 - , - , role of crosslinking agent 325-327 - , crosslinking density 348, 366 - , - , a m i n o resins, of 389 - , - , epoxy resins, of 348 - , crosslinking degree effect 405 Cross-over t e m p e r a t u r e 408 Crystallinity 461 Crystallization of chain molecules

621 kinks 295 polyisoprene (cis 1, 4) Curing of epoxy resins 359 ff Damage cummulation 479 Debonding 569,575 interfacial 569 Defect - , cohesion 481 - , composite structure 221 ends of microfibrils 14 Deformability of ABS copolymers 278-280 Deformation 143,144 - , conformational 220 - , creep 210 matrix microvolumes, of 509 - , mechanism of 143,144 necking and drawing, with 512 - , plastic 217 reversible 217 Degree of - , crosslinking 373 - , dispersion 243, 247, 612, 613 foaming 226-232 Delamination 301 Dielectric constant 75-77, 80 Dielectrics polarization/depolarization 363 Diethyl phthalate 31 Differential scanning calorimetry (DSC) 459,460 Diffuse reflectance spectroscopy (DRS) 452 Dilatometry 7 Discharge thermally stimulated 351 Discontinuous fibres 183-187,190 Dispersant 243,244, 246-248 Dispersion of fillers 607, 608-612 Dispersion aids 159 - , grafted cellulose 155,160 - , hydrolysis of cellulose 155 - , - , disintegration of fibres 155,162 - , - , improved dispersion 162 - , - , improved modulus 162 - , - , microfibrils 153,162,164 Dispersion degree 243, 247, 612, 613 Distribution of relaxation times, of 365 Drying 335 Ductile fracture 508 Dynamic mechanical properties 348, 450, 451, 460, 539, 541, 545 Dynamic modulus 453-456 Effective notch length 238 Effect of filler shape 19-32 Elastic bulk modulus 106,113, 400 Elastic-viscoelastic problems 107 Elastomer - , butadiene-styrene 275-278

- , EPDM, modified composite 213, 218 - , interphase 373, 375, 376,, 378 Electrical properties of metal filled composites 81 Electric field 176-181 Electrokinetic method 417, 420 Electroosmotic method 414, 427 Electrophoresis 417,427 Electrophysical properties 71 - , charge carriers 77, 79 - , concentration threshold 72 - , conductivity 72 - , conductivity in periodic fields 77 - , dielectric 76 - , dielectric constant 75, 77, 80 - , insulating layer 76,79 - , local conductivity 74 - , metallic conductivity 85 - , ohmic contact 80 - , percolation threshold 72, 73,75, 76 - , random potential barrier 79 - , resistance 76, 84 - , thermal fluctuations 76, 77, 82 - , transfer of charges 76 - , tunneling 76, 77, 82 Elongation 156-162 - , at break 168, 307, 558 - , of fibres 54, 55 Embryonic spherulites 34 Energy dissipation 217 Energy consumption 217 Epitaxy - , epitaxial growth 7 Epoxy network 373 Epoxy resins - , anhydride cured 349 - , blends with phenolic microspheres 382 - , - , mechanical properties 384, 385 - , - , morphology 351, 387 - , catalyst cured 349 - , crosslinked density 348 - , domain structure 348 - , electron spectroscopy 348 - , impact properties 348 - , inhomogeneities 348 - , low temperature relaxation 347 - , nodules 348 - , relaxation mechanisms 347 - , SAXS 348 - , tensile strength 348 Equal load sharing rule (ELS) 52 Ethylene-vinyl acetate 397, 401-406 Expansion coefficient, thermal 407 Extensional rigidity 63 Failure 105,116,118, 452 - , adhesion strength 497, 499, 501-503, 505, 506

622 cohesion defects 481 - , cohesion failure of fibre 497, 505 complex fatigue process 480 cracks and microcracks 480 - , crazes 484 damage cummulation 479 failure process velocity 481 hybrid composite 489, 494 - , macrocrack and fracture formation 481 - , multiple 64 - , pattern 63 prestrained fibre joints 497-505 - , process 44, 64, 491, 494, 511 - , propagation of cracks 483 - , reliability of structures 479 - , residual stresses 497-501, 504,505 - , roving strength 492, 493 - , rupture elongation 488, 489,492 - , shear strength 497-506 shrinkage 497,499, 500, 505 strain 58 - , strand-strength 489, 490 stress concentration on fibre surface 485 - , tensile strain 501 - , tensile strength 488-490 - , Weibull P-parameter 490-492 - , yarn strength 487-490 Fibre 247, 248 - , carbon 37, 38, 41 - , critical length 45, 46,187, 594 - , discontinuous 183,186,187 glass 58,183,184, 218, 507-512, 575, 576 - , hardwood 252 - , - , chemithermomechanical pulp preparation 252 pulp grafting 253 - , high elongation 54, 55 length 183,184 - , low elongation 54, 55 - , reinforcement 183 - , slag-basalt fibres 284 slag fibres 284 Fibre orientation - , determination 589 laminate microstructure 589 - , programme CAFOD, application 589, 591-594 - , - , angular distribution 589 - , - , number and size of layers 589 - , - , orientation parameter 589, 590 transition zones 589 - , - ,