Advanced Routes for Polymer Toughening

Polymer Science Library 10 ADVANCED ROUTES FOR POLYMER TOUGHENING Polymer Science Library Edited by A.D. Jenkins Uni...

Author: E. Martuscelli | P. Musto | G. Ragosta

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Polymer Science Library 10

ADVANCED ROUTES FOR POLYMER TOUGHENING

Polymer Science Library Edited by A.D. Jenkins University of Sussex, The School of Molecular Sciences, Falmer, Brighton BN1 9QJ, England 1. K. Murakami and K. Ono, Chemorheology of Polymers 2. M. Bohdaneck~ and J. KovdL Viscosity of Polymer Solutions 3. J. Wypych, Polyvinyl Chloride Degradation 4. J. Wypych, Polyvinyl Chloride Stabilization 5. P. Kratochvfl, Classical Light Scattering from Polymer Solutions 6. J. Bartoh and E. Borsig, Complexes in Free-Radical Polymerization 7. Yu.S. Lipatov, Colloid Chemistry of Polymers 8. F.J. Baltd-Calleja and C.G. Vonk, X-Ray Scattering of Synthetic Polymers 9. K. Kamide, Thermodynamics of Polymer Solutions

Polymer Science Library 10

ADVANCED ROUTES FOR POLYMER TOUGHENING

Edited by

E. Martuscelli P. Musto G. Ragosta National Research Council of Baly, Institute of Research and Technology of Plastic Materials, Arco Felice (Naples), Italy

1996 ELSEVIER Amsterdam

- Lausanne

- New

York - Oxford

- Shannon

- Tokyo

ELSEVIER SCIENCE B.V. Sara Burgerhartstraat 25 P.O. Box 211, 1000 AE Amsterdam, The Netherlands

ISBN: 0-444-81960-6 (vol. 10) ISBN: 0-444-41832-6 (Series) 9 1996 Elsevier Science B.V. All rights reserved. No part of this publication may be reproduced, stored in a retrieval system or transmitted in any form or by any means, electronic, mechanical, photocopying, recording or otherwise, without the prior written permission of the publisher, Elsevier Science B.V., Copyright & Permissions Department, P.O. Box 521, 1000 AM Amsterdam, The Netherlands. Special regulations for readers in the USA. This publication has been registered with the Copyright Clearance Center Inc.(CCC), 222 Rosewood Drive Danvers, MA 01923. Information can be obtained from the CCC about conditions under which photocopies of parts of this publication may be made in the U.S.A. All other copyright questions, including photocopying outside of the USA, should be referred to the copyfight owner, Elsevier Science B.V., unless otherwise specified. No responsibility is assumed by the publisher for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions or ideas contained in the material herein. This book is printed on acid-free paper. Printed in The Netherlands

CONTENTS Preface ...................................................................................................... vii

Introduction ................................................................................................

1

E. Martuscelli, P. Musto, G. Ragosta

P A R T I: T O U G H E N E D

THERMOSETS

Chapter 1. Epoxy Resins .......................................................................... 11

E. Martuscelli, P. Musto, G. Ragosta, G. Scarinzi

Chapter 2. Unsaturated Polyester Resins ................................................. 61


Chapter 3. Thermosetting Polyimides .................................................... 121


P A R T II: T O U G H E N E D

THERMOPLASTICS

Chapter 4. Nucleation Processes in Toughened Plastics ......................... 157

A. Galesh, Z. Barwzak, E. Martuscelli Chapter 5. Isotactic Polypropylene Based B l e n d s .................................. 243

L. D'Orazio, C. Mancarella, E. Martuscelli, G. Sticotti

Chapter 6. Polyamide 6 / E t h y l e n e - c o - V i n y l a c e t a t e B l e n d s ...................... 2 8 9 a Model System of Thermoplastic/Elastomer Pairs

L. D'Orazio, C. Mancarella, E. Martuscelli

vi

Chapter 7. Blends Polyamide 6/Functionalized Rubber . . . . . . . . . . . . . . . . . . . . . . . . .

335

R. Greco, M. Malinconico, E. Martuscelli, G. Ragosta

Chapter 8. PMMA/Rubber Blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

439

P. Laurienzo, M. Malinconico, E. Martuscelli, G. Ragosta, M. G. Volpe

Chapter 9. Polycarbonate Toughening by ABS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

469

R. Greco

Chapter 10. Rubber Modification of Biodegradable Polymers .............. 527 M. Avella, B. Immirzi, M. Malinconico, E. Martuscelli, M. G. Volpe, M. Canetti, P. Sadocco, A. Seres

vii

PREFACE The problem of polymer toughening elnerged more than twenty years ago and since then has experienced a steady increase in interest aald in the number of researchers all over the world dedicating their efforts to its solution. Although many aspects remain to be clarified and many questions still wait for a definitive answer, nowadays toughened plastics hold a consistent and well defined position ha the polymer market and successfully withstand the competition of advanced engineering plastics whose R&D costs are often prohibitive. hi such a complex and rapidly evolving scenario, we were invited by Professor A. D. Jenkins, Editor in chief of the Polymer Science Library edited by Elsevier Science, to produce a book on this topic. At a first glance we were hltrigued by this opportunity but soon realized how extensive and comprehensive the existing literature is, especially when considering a number of highly appreciated and widely diffused books (see, for instance, the references cited in the introduction). Therefore we decided not to attempt a comprehensive coverage of all the aspects of the subject matter, but rather to concentrate the attention on a limited number of systems which may be regarded as typologies of toughened plastics. From an extensive treatment of these systems we felt it possible to develop general concepts whose importance we learnt to appreciate more mad more as our research activity in the field has proceeded. These concepts, which can never be overestimated in the design mid formulation of tough polylner blends, include the role of the interface in lnultieomponent systems, the chemical reactivity of the blend components, the mode mad state of dispersion of the second phase, the crystalli~fity mad crystallization conditions, the glass transition temperature etc. The present book therefore mainly reports on the long term activity of the Institute of Research and Teclmology of Plastic Materials in the area, which started in the early seventies with the chemical modification of ethylene-propylene copolymers to be used as tougheners of Polyamide-6. The hlteresthlg results obtained prompted us to investigate further such a system both from a fundamental

viii and from an applicative point of view, and from this initial and highly rewarding experience we gained an approach which we have never abandoned along the years, i.e. to make fundamental research on systems of high technological interest. In this respect an essential part of our activity has always been represented by our continuous contact with research teams operating in the industrial world who gave us numerous inputs aald faced us with the problems of the demanding market of innovative materials. We hope that the present contribution can give to the reader the feeling of how intellectually sthnulating aald challenging the field of polymer toughening can be, so as to encourage him to spend his research efforts in such an area. Finally we wish to express our deep gratitude to all the authors who have contributed to this volume and to the many others behind them, each of one can be found in the extended literature cited. A special thank goes to Mr. Giuseppe Narciso, who gave us an invaluable help in the hard task of editing the manuscript.

Ezio Martuscelli PeUegrino Musto Giuseppe Ragosta

National Research Council of Italy; Institute of Research and Teclmology of Plastic Materials Arco Felice (Naples), Italy.

INTRODUCTION


National Research Council of Italy, Institute of Research and Technology of Plastic Materials, 80072 Arco Felice (Na), ITALY

Toughened polymers represent a large area of scientific and technological concern. In fact, with the gradual penetration of plastics in areas traditionally dominated by metals and ceramics, new polymeric materials, both thermoplastics and thermoset resins, have been developed for increasingly demanding applications. These materials are able to provide the right combination of lightness and mechanical performance over a wide range of temperatures. In mm~y instmaces a good balance between stiffness and tougheness is required but most of the as sinthesized polymers which exlfibit adequate rigidity are characterized by brittleness and low resistance to crack propagation. Therefore the increase of the intrinsic tougheness of otherwise brittle materials has stimulated a huge mnoum of research efforts both in the academic mid in the industrial world [ 1, 2, 3]. Two different approaches have emerged, both with their own potential and weakness. One is to sinthesize new omo- or copolymers, based on novel monomers as in the case of polycarbonates, polysulphones and polyether-ketones. The second approach consists in modifying existing polymers through the addition of a second polymeric component, a route generally referred to as "blending". Such a method presents the distinctive advantage of being, in general, more economically attractive, since the development of new sinthetic methods is a long and costly process [4].

An excellent example of the blending approach is provided by the rubber toughening in which a small amount of rubber, typically between 5 and 20 % by weight, is incorporated as a disperse phase into a rigid plastic matrix. The resulting blend is characterized by a considerably higher fracture tougheness than the parent polymer; there is an inevitable reduction in the modulus and tensile strength but these losses are far outweighed by the improvement in fracture tougheness. This approach has proven to be very successful and a wide variety of plastics toughened in this way are now commercially available. Among the best known examples are high impact polystyrene (HIPS), and polyvinylchloride (PVC); other plastics which have been toughened using this teclmology include polymethylmethacrylate (PMMA), polypropylene (PP), polycarbonate (PC), Nylons and, most recently, thermosetting resins such as epoxies, polyimides and unsaturated polyesters [5, 6]. In the case of rubber toughened thermoplastics an essential condition to achieve satisfactory results is that the rubbery phase must be finely and homogeneously dispersed within the matrix; furthermore the rubbery particles must be adequately bonded to the matrix. To this end the rubber must posses a solubility parameter sufficiently different from that of the matrix polymer to ensure a fine second phase dispersion but close enough to promote adequate adhesion of the particles to the matrix. Such stringent requirements strongly limit the choice of possible rubber tougheners for a given polymer matrix. To overcome this limitation the concept of compatibilizing agents in the form of block or graft copolymers has been developed and succesfully applied to a wide number of actual cases [7, 8]. Essentially, a suitable

block or graft copolymer whose segments are

chemically equivalent to the blend components is added. Atter the blending process such a copolymer locates preferentially at the interphase thus promoting a better dispersion of the second component and an improved adhesion among the phases. The behaviour of small amounts of compatibilizer in an immiscible blend has been described as that of a classic emulsifying agent, similar to the soap molecules at an oil-water interface [9]. The success of the use of block or graft copolymers (and in some instances also random copolymers) as compatibilizers accounts for many of

the large number of commercially available blends, e.g. HIPS and ABS. The renewed interest towards the reactive melt processing, i.e. reactive extrusion (REX) and reactive injection moulding (RIM), is also due to the recent achievements in the development of copolymer compatibilizers. Moreover, when the thermoplastic matrix is able to crystallize, further factors must be taken into account such as the structure and size of spherulites and lamellar crystals, the spherulite grow rate and the nucleation process, which are all affected by the presence of the dispersed sott phase [ 10-13]. The toughening of thermosetting resins poses different, yet equally challenging issues. Generally, a critical step towards the preparation of a toughened them~osetting blend is to start from a single-phase homogeneous reactants mixture prior to the curing process. For this reason the rubbers usually employed are low molecular weight liquids and are miscible enough to dissolve in the resin prepolymer. However the elastomeric phase separates out during the curing process, giving rise to fine and homogeneous dispersion of the second component in the resin matrix. The phase separation process is of paramount importance in these systems and depends on both kinetic and thermodynamic factors. The understanding of these factors and the ability to control them to obtain the desired morphology of the materials represent one of the main goals of the research efforts in this area. As for thermoplastics, also in the case of thermosetting blends, a certain degree of chemical interaction between the resin and the rubber modifier is required to improve the interfacial adhesion and hence to achieve an effective toughening. Such reactions generally involve the end-groups of the rubber modifier whose functionality is adjusted according to the chemical nature of the matrix resin. An example of an extensively investigated system of this type is provided by the blend composed by a bifunctional epoxy resin and a carboxyl-terminated acrylonitrile-butadiene copolymer (CTBN). This system is particularly versatile since the solubility of the CTBN modifier in the uncured epoxy resin can be adjusted by varying the acrylonitrile content in the copolymer. However the CTBN precipitates out from solution during curing and, at the same time, a chemical

reaction occurs among the carboxyl end-groups of the rubber and the oxirane rings of the matrix. The rubber toughening approach of thermosets fails when high temperature requirements becomes more and more critical. In fact the reduced themml stability of unsaturated, low molecular weight rubbers, renders such blends unsuitable for high temperature applications. Moreover the toughening efficiency of elastomeric modifiers gradually decreases as the cross-linking density of the matrix increases. This is because, as the thermoset resin is cross-linked more tightly, its capability to be plastically deformed is strongly reduced and the dispersed rubber particles are no longer able to induce energy dissipation processes in the matrix. Both of these issues are particularly relevant in the case of the latest, high teclmology termosets like tetrafunctional epoxies and bismaleimide resins which find application in the aerospace and in fl~e electronic industries. For these materials a novel approach has recently emerged, which consists in the fommlation of blends with tough, ductile and thermally stable engineering thermoplastics. Examples of thennoplastics successfully employed in such blend systems are provided by polycarbonates, polyetheresulphones, polyarilsulphones and polyefllerimides. Concurrently with the development of novel tougheners and of more sofisticated technologies to produce multicomponent polymer blends with balanced end-properties, a large amount of efforts has been spent to elucidate the mechanisms of fracture in these complex systems. This in an attempt to be able to control the very many factors which play a role in the fracture behaviour of toughened plastics. Alter more than twenty years of extensive research in this area we may say that we are still far from a complete understanding of the whole phenomenon but very significant advances have been achieved especially in the case of blends based on rubber modifiers. In fact it is now well established that rubber particles with low moduli act as stress concentrators in both thennoplastic and themloset resins, enhancing shear yielding and/or crazing in the matrix, dependhlg on its molecular architecture. In particular in the case of themlosets the

crazing mechanism does not operate, while one important process is the initiation and growth of multiple localized shear yield deformations in the matrix. In addition a cavitation process occurring either in the rubber particles or at the particle-matrix interface often plays a key role. Once formed, these voids grow and so dissipate energy; at the same time they lower the stress required to initiate shear yielding in the matrix thus promoting more extensive plastic deformation [ 14]. hi the present book the theme of polymer toughening is covered in its very diverse aspects. Each chapter deals with a particular material or class of materials and is aimed at giving an overview of the problems encountered and of the proposed solutions. In particular the first part is devoted to thermosetting resins with three chapters dealing with epoxy resins, unsaturated polyesters and bismaleimide resins. The toughening of highly cross-linked epoxies by blending with engineering thennoplastics is covered in Chapter 1. Two different approaches are presented, depending on the nature of the thermoplastic modifier. With bisphenol-A polycarbonate a reactive blending process was developed, by which polycarbonate chains were chemically h~corporated within the epoxy network. A non-reactive blending process was carried out when a polyetherimide was used as toughened. In both the cases satisfactory results were achieved. The second chapter describes the toughening of unsaturated polyester resins by reactive liquid rubbers. Commercially available elastomers were chemically modified in order to enhance their compatibility and reactivity towards the polyester resin. These modified rubbers were molecularly characterized and their chemical

interactions

with the

matrix

during

curing

were

investigated

spectroscopically. The morphology and the fracture properties were found to be strongly dependent on the type of rubber modifier and on the chemical nature of its end-groups. The third chapter on thermosets deals with the toughening of polyimides. For this class of materials the problem has been approached either by adding a themloplastic second component (a polyetherimide), or a reactive liquid rubber. Better results were achieved with the polyetherimide, despite a number of

processing problems encountered along the way. This confirms the reduced efficiency of rubber modifiers in the toughening of very highly cross-linked themaosets. The second part of the book, concerned with thermoplastics, is opened by a chapter on the nucleation processes occurring in toughened semicrystalline plastics. This is a subject of great theoretical and practical interest, and a comprehensive account is given of the most recent advances in the interpretation of such a complex phenomenon. The role played by file molecular and microstructural parameters of the components on the toughening of blends of isotactic polypropylene (iPP) and ethylene-propylene (EPR) copolymers is treated in detail in Chapter 5. Here is shown that the desired fmal properties can be imparted by an appropriate choice of molecular mass and molecular mass distribution, constitution and tacticity of the blend components. A further critical factor is represented by the crystallization conditions through which it is possible to optimize the mode and state of dispersion of the rubbery phase in the semicrystalline matrix. Chapters 6 and 7 both deal with blends based on polyamide-6 (PA6). The fonner treats a blend system whose minor components are ethylene-vinylacetate (EVA) copolymers. The latter chapter covers PA6 blends in which the modifier is a functionalized EPR copolymer. These systems have received a great deal of attention in our Institute and, historically, represent the starting point of our continuing involvement in the field of polylner toughening. Particularly interesting results were obtained with an EPR elastomer functionalized by hlsertion of succinic alfllydride groups along its backbone. Such a modified rubber was used either as interfacial agent in PA6/EPR blends or directly as toughener for the PA6 matrix. Two different routes were followed to prepare the blends: by melt mixing the components in a Brabender-like apparatus or concurrently with the hydrolytic polymerization of e-caprolactam. The results obtained by these two approaches are critically reviewed. In chapter 8 the problelns involved with the toughening of amorphous thermoplastic polymers such as polymethylmethacrylate (PMMA) mid polystyrene

(PS) are discussed. In particular a novel method for the PMMA toughening is presented, which appears simpler to be implemented for large scale applications then those already established. Such a method consists in the dissolution of the rubber modifier (EVA copolymers) in file acrylic monomer which is subsequently radically polymerized. Satisfactory results in terms of tougheness were obtained even with limited amounts of the second component in the blend. Chapter 9 deals with the toughening of polycarbonate (PC) by ABS copolymers. The chapter starts with a critical review of vast amount of literature data, stressing the points which deserve further investigation. Subsequently the results obtained by the author's research group on aspects such as processability, miscibility and thermal behaviour of these blend are presented and discussed. Finally the last chapter is concerned with a class of biomaterials like poly (13hydroxybutyrate) (PHB) and his copolymers produced by bacterial synthesis. These polymers are receiving increasing attention due to their biocompatibility and biodegradability and the improvement of their mechanical properties is becoming a problem of considerable technological interest. Recent results obtained using methods based on bulk or suspension pol3qnerization of the modifier are discussed. References

1.

C.B. Bucknall, "Toughened Plastics", Appl. Sci. Pub., London, 1977.

2.

A.J. Kinloch, R. J. Young, "Fracture Behaviour of Polymers", Appl. Sci. Pub., London, 1983.

3.

J. M. Margolis, "Advanced Thermoset Composites", Van Nostrand Reinhold Co., New York, N. Y., 1986.

4.

M . J . Folkes, P. S. Hope, "Polymer Blends and Alloys", Blackie Academic & Professional, London, 1993.

5.

C. K. Riew, A. J. Kinloch, "Toughened Plastics I: Science and Engineering" Advances in Chemistry Series, 233, ACS, Washington, D.C., 1993.

~

C. K. Riew, "Rubber Toughened Plastics", Advances in Chemistry Series, 222, ACS, Washington, DC, 1989.

.

E. Martuscelli, R. Palumbo, M. Kryszewski, "Polymer Blends", Vol. I, Plenum Press, New York, 1980.

.

A. Galeski, E. Martuscelli, M. Kryszewski, "Polymer Blends", Vol. II, Plenum Press, New York, 1984.

.

D.R. Paul, "Polymer Blends ", Vol. 2, D.R. Paul and S. Newman Eds., Academic Press, London, 1978.

10. E. Martuscelli, "Rubber Modification of Polymers: Phase Structure, Crystallization, Processing and Properties" in "Thermoplastic Elastomers from Rubber-Plastic Blends", S. K. De, A. K. Bhowmick Eds., Ellis Horwood, New York, 1990. 11. E. Martuscelli, "Structure and Properties of Polypropylene-Elastomer Blends" hi "Polypropylene: Structure, Blends and Composites", J. Karger-Kocsis Ed. Chapman and Hall, London, 1995. 12. Z. Bartzak, E. Martuscelli, A. Galeski, "Primary Spherulite Nucleation in Polypropylene-based Blends and Copolymers" in "Polypropylene: Structure, Blends and Composites", J. Karger-Kocsis Ed. Chapman and Hall, London, 1995. 13. E.

Martuscelli, "Relationships Between Morphology, Structure,

Composition and Properties in Isotactic Polypropylene Based Blends" in "Polymer Blends and Mixtures", D. J. Walsh, J. S. Higgins, A. Macolmachie, NATO ASI Series, 1984. 14. A. C. Roulin-Moloney, "Fractography and Failure Mechanisms of Polymers and Composites", Elsevier Appl. Sci., London, 1988.

PART 1 TOUGHENED T H E R M O S E T S

This Page Intentionally Left Blank

11 CHAPTER 1

EPOXY RESINS E. Martuscelli, P. Musto, G. Ragosta, G. Scarinzi National Research Council of Italy, Institute of Research and Teclmology of Plastic Materials, 80072 Arco Felice (Na) ITALY.

1. Introduction

Recently epoxy resins with high crosslinking density have received considerable attention owing to their high temperature performances [ 1-3]. These high technology materials find applications in the aerospace industry as well as in other structural applications where the stiffness, the creep resistance for extended periods of time mad the thermal stability are essential [4]. The chemical formulas of the commercially available multifunctional epoxies are reported below:

/ok

A /~/c.=--CH---CH=

CH2--OH--CH2~~N \

TGAP:TriglycidylepoxJdebasedon aminophenol

/o,,,

A

CH2"--CH--CH2\ ~ /~---~ /CH2~CH~CH 2 N---~( ) )---C H2----'~~ ) )'----N C~2?H--CH / ~ / \'~--J/ \CH2__CH__CH2

V

o

TGDDM:TetraglycidilepoxJdebasedon DDM

12

O

/ 0\

O

/\

O-CH2-CH-CH2

O-CH2-CH-CI- ~

/\

O-CH2-CH-CH =

0

Novolac resins: Average value for n: DEN 438 = 1.6 DEN 439 = 1.8

Those resins are crosslinkcd with a number of different hardeners, the more common being reported below:

H2N---( (

) )-----.S.~

(~

))----NH 2

DDS: 4, 4' diarnino diphenyl sulfone

H 2 ~ C H 2 ~ N H 2 DDM: 4, 4' diarnino diphenyl methane

MNA: methylnadic anhydride

When those resins are fully cured, the resulting molecular network exhibits very high values of glass transition temperatures, Tg,(>250~

compared to the

conventional bifunctional epoxies. The stiffness of these materials is maintained up to temperatures approaching the Tg.However, as for other thermosets, the fracture

13 tougheness is extremely low, limiting the application of these materials to situations where the stress is relatively low and preferably static. The well developed approach of using a reactive liquid rubber as a toughening agent, which has meet with considerable success in the case of bifimctional epoxies [5-7], does not provide significant improvements for highly crosslinked resins. The reason why this is so has been dearly demonstrated by Yee and Pearson [8,9]. They tested a series of resins made from epoxies with different starting molecular weight. Their results, reported in Fig. 1, show that, in the absence of rubber, an increase in the molecular weight of the starting resin, which corresponds to a decrease of the cross linking density, has a very limited effect on the fracture toughness. On the contrary, when a rubber is added, the resins with higher cross-link density exhibit a very small improvement in toughness, which increases substantially by decreasing the cross-link density.

---

7

E #

6-

13.. w

I

I

I

9 Rubber modified o unmodified

5-

(/)

In a)

r

J~ o~ =3 0

I-(D

4-

3-

2-

IL_

::3

(J

1 -

(o

IL

u.

0 0

I

I

I

1000

2000

3000

Epoxy

Equivalent

4000

Weight

Figure 1. Fracture tougheness values for a series of DGEBA-DDS epoxies. (After Yee and Pearson).

14 From the point of view of the deformation mechanism this effect can bc cxplaine,d considering that, as the resin is cross-linked more tightly, its capability to bc plastically dcformexi is strongly rexiuce~ and the dispcrsexl robber particles arc no longer able to induce energy dissipation mechanisms such as shear yielding and/or cavitation [10-14]. Therefore the higher is the cross-linking density of the matrix, the lower its toughcnability. Recently a novel approach to toughening of epoxies has emerged, which has the potential of solving the above discussed problems. It consists in the formulation of blends with tough, ductile and thermally stable cngin~ring thermoplastics [15-21]. This altcmativc approach has the additional advantage that other desirable properties of the matrix, such as modulus, yield stress and glass transition temperature, arc not adversely affcctexi by the addition of the modifier.

80

I

i

3.5

I

o Yield Stress m

9 Modulus

m

70

t~

- 3.o

5

E t_

o

9 60-

r,,, .,,.,.==

t,,ol

v

"0

- 2.5

l

50-

4O 0

i

I

10

20

Rubber Content (phr}

Figure 2. Mechanical properties of a CTBN-modified epoxy.

2.0

30

,.,. Q "10

15 On the contrary Fig. 2 demonstrates a considerable decrease of these properties when a reactive liquid rubber is employed as second component [22]. The thermoplastic component can be either incorporated within the thermosetting network by means of a suitable reactive blending process or can be simply mixed with the matrix, giving rise, in most cases, to a completely phase separated system after the curing process. The thermoplastic modifiers cited in the literature are reported below:

c.3

I~

PC: Poly(bisphenoI-Acarbonate)

r,

o

!

] ~n

PEI:a thermoplasticPoly(etherimide)

PSU:Poly(sulfone)

PESU:Poly(ethersulfone)

16 The PSU, PESU based systems have been investigated by several authors [2334]. Some of these systems were found to phase separate after the curing process to give modifier's particles of about 0.5 pm in diameter. Unfortunately the resulting improvement in fracture toughness was modest. These results prompted to investigate other types of thermoplastics, and in the present chapter the systems DEN 438/PC and TGDDM/PEI are discussed in detail.

2. DEN 435/PC system 2.1 Molecular characterization and curing kinetics. A critical step toward the preparation of a successful thermosetting blend is to start from a single-phase homogeneous reactant mixture prior to the curing process. In our case this was achieved by a reactive blending process in which the PC was dissolved at high temperature (220~

for extended periods of time (typically 3 h) in

the uncured resin. At the end of this procedure, a clear, homogeneous solution was obtained for the whole composition range investigated. Then the temperature of the mixture was allowed to decrease to 80~ and the hardener MNA and the accelerator BDMA were added. The mixture was then poured into a glass mold, and the curing process was carried out at 120~ for 20h. A final posteuring step was performed for 5 h at 200~

After this protocol, a visually transparent sheet was recovered. The

various investigated blend compositions, together with their codes are reported in Table 1. The DEN 438/PC mixtures, prior to the addition of the hardener, were examined by differential scanning calorimetry (DSC) and by Fourier Transform Infrared Spectroscopy (FTIR) to insure the miscibility of the system. DSC measurements (Fig. 3) showed that all the mixtures investigated exhibit single glass transition temperatures, Tg, intermediate between those of the pure components, which, as expected, increase with increasing the PC content in the mixture.

17 Table 1.

Codes and Compositions of the Epoxy~PC Blends Code

Epoxy

MNA

Epoxy +

(%)

(%)

B0

49.5

50.5

100

B4

47.5

48.5

96

4

B8

45.0

47.0

92

8

B 10

44.5

45.5

90

10

B 12

43.5

44.5

88

12

B15

41.5

43.5

85

15

B20

39.6

40.4

80

20

MNA

PC

(%)

PC

_66134

9zla

-20

I

l

i

~,

20

60

100

140

T oC

I 180

Figure 3. DSC traces of uncured epoxy resin and of uncured epoxy~PC mixtures. Compositions as indicated.

18 In Fig. 4 the Tg values are reported as a function of composition. This behavior indicates that the high-temperature dissolution process produces a single-phase, homogeneous system over the entire composition range investigated.

310

I

I

I

300 -

lit

I---

290

280 0

I

I

I

0.1

0.2

0.3

0.4

PC Weight Fraction

Figure 4. The glass transiaon temperature, T~ as a fimction of composition for the Epoxy~PC mixtures investigated. Fig. 5 gives the FTIR transmission spectrum in the range 4000-400 cm -1 of a film of the uncured resin cast from CH2CI 2. Characteristic absorptions due to hydroxyl groups are observed in the 37003200 cm -1 range. The C-H stretching region (3200 - 2880 cm -1) is complex and highly overlapped while in the lower frequency range (1600-400 cm -1 ) a number of better resolved peaks are detected. For some of them a well-defined baseline can be identified, thus affording, after proper assignment, their use for analytical purposes.

19 The complexity of the resin's molecular structure makes complete assignment based on normal-coordinate analysis an extremely difficult task. Thus the peaks' assignments rely heavily on the work of Antoon [35] as well as on those of Bellamy [36] and of Colthup et al. [37].

.4-

_j 0-~

,

4000

'~

aooo

20'00 Wavenumbers

x6oo

( e r a - - 1)

Figure 5. FHR transmission spectrum o f uncured epoxy resin at 4000-400 cm -1. The spectrum was obtained on a film cast from CH2CI~

1.5

0 4000

30'00

2obo

Wavenumbers (era-- I)

~o'oo

Figure 6. bTIR transmission spectrum o f polycarbonate at 4000-400 cm -I. The spectrum was obtained on a thin film cast from CH2CI2.

20 In Fg. 6

is reported the transmission FTIR spectrum of PC: the carbonyl

stretclfing vibration produces a strong, well-resolved peak at 1775 cm-l, while the O-C-O stretching mode gives rise to a very intense and complex multiplet with maxima at 1227, 1193, and 1163 cm-1. In the spectrum of a 66/34 epoxy/PC mixture the C=O peak of PC is shifted by 3 cm-l with respect to the peak position detected in pure amorphous PC. This effect is evidenced in Fig. 7A where the spectra of the mixture and of pure PC are compared in the 1850-1700 cm-l range. Fig. 7B shows the result of spectral subtraction performed on the mixture spectrum using pure PC as the reference: the derivative-type feature characteristic of band shift in the reference is evident.

A / \,\

77J~-/~

\Ix.~_.~

,,Q

18bo Waven.umbene

175o (cm-- I )

~

"

tsoo Wavenumbera

t~5o (era-- I )

Figure 7. A) transmission spectra of an 80/20 epoxy~PC mixture o f pure PC m the 2000-1700 cm "1 range; B) subtraction spectrum in the 20001700 cm -I range obtained from the 80/20 epoxy/PC mixture using pure PC as reference.

Previous studies on the vibrational behaviour of PC in solution [38] demonstrated analogous shifts of the C=O peak in chlorinated solvents like CH2CI2 and CHCI3; in particular the shift was increased by increasing the protonating

21 properties of the solvent (-3.0 cm -I for CH2CI 2 and -4.5 cm-I for CHCI3 in a 1% w/w solution). The above effects were interpreted in terms of hydrogen bonding interactions between carbonyl oxygen of the polymer and hydrogen of the solvent. An analogous type of molecular interaction can be assumed to account for the effect observed in the epoxy/PC mixture. In particular, in this case hydrogen bonding may involve the PC carbonyls as proton acceptors and the methylene or methine groups bonded to the electron withdrawing oxygens, as proton donors. The fact that in the mixture spectrum we observe a single, highly symmetrical Vc_o absorption (see Fig. 7A) indicates that most of the PC carbonyls are involved in the above interactions. This in turn implies that the PC is molecularly dispersed in the epoxy matrix, thus confirlning the DSC results which indicated that the epoxy/PC mixture is a single-phase, miscible system. A deeper analysis of the uncured epoxy/PC mixture has been performed in order to investigate whether chemical interactions among the components occur in addition to the previously discussed physical interactions during the dissolution process.For this purpose the PC component was selectively extracted from the system: the mixture was first completely dissolved in a common solvent (CH2C12) and then the PC component was precipitated from solution using acetone, which is a solvent for the epoxy resin and a non solvent for PC. An epoxy/PC mixture of the same composition, but obtained by direct dissolution of the components in CH2CI2 was used as a reference system. The results of the quantitative extraction are summarized in Table 2. It is noted that, while in the reference system the PC is almost completely recovered (ca. 95%), in the mixture obtained by high-temperature dissolution of PC in the resin, only about 50% of PC can be extracted. Also of interest are the GPC data of the various recovered fractions reported in the last three columns of Table IV. It is noted that the PC fraction recovered from the reference system displays a slight increase in both M w and M~ while the molecular weight distribution is narrower with respect to the starting PC. This is likely due to the fact that, during acetone extraction, the lower molecular weight fractions of PC do not precipitate; in fact we recovered only 95% of the total extractable fraction.

22 Table

2.

Results of Quanatative Extraction on Epoxy~PCMixtures and GPC data m

Epoxy

PC

extracted

M . 10-4

M . 10-4

M , / M~

(%)

(%)

Pc (%)

0

100

-

1.2

3.1

2.6

73.5

26.5 b

95

2.4

3.9

1.6

73.5

26.5 c

47.1

1.3

1.9

1.5

aon the total amount of extractable PC. bmixture prepared by direct dissolution in CH2CI2. cmixture prepared by high-temperature dissolution. However the most relevant observation is that the PC fraction recovered from the epoxy/PC mixture obtained at high temperature, shows a marked decrease in the molecular weight moments, while the polydispersity remains scarcely affected. In particular both and M w are about half of the values detected in the reference system. This indicates that, during the high temperature dissolution process, the PC undergoes chain-scission reactions. Moreover the observation that M~ and M~ decrease by about half may be interpreted by assuming a random chain-scission process with no preferential sites along the PC backbone. These processes also account for the lower amount of extractable PC in the case of the epoxy/PC mixture. In Fig. 8 is reported the FTIR transmission spectrum of the PC fraction recovered from the epoxy/PC mixture. The measurement was performed on a thin film cast from CH2CI 2. Comparison with the pure PC spectrum (see insets) evidences the occurrence of a broad absorption centred at about 3528 cm -1, together with a lowintensity peak at 916 cm-1. The first characteristically broad contribution can be readily ascribed to stretching vibrations of self-associated hydroxyl groups. The band at 916 cm"l is attributed to a ring mode characteristic of epoxy groups. Further information on the VoH absorption can be gamed by eliminating the interference of hydrogen-bonding interactions, which, as is well known, produce extensive band broadening and make the spectral region poorly resolved [36,37].

23

0.58

~

0.40

' "

.Q

'

I

9 A I

'

4000

B

2500

i

O

.Q 0.23

0.05 !

i

3622.2

i'

i

1 ....

i"

1

2866.7 211 I. I 1355.6 Wavenumber ( c m -I )

600.00

Figure 8. FTIR transmission spectrum of the PC fraction recovered from the 80/20 epoxy~PC mixture. The spectrum was measured on a thin film cast from CH2CI~ The insets compare this spectrum with that of the pure PC in two different frequency ranges.

The spectra of dilute solutions of this PC fraction and of pure PC in a low polarity solvent (CH:CI2) are reported in Fig. 9, traces A and B, respectively. A sharp peak at 3583 cm-1, not present in the starting polymer, is evident in the PC recovered fractions; this band is partially overlapped with a PC absorption (a Vc_o overtone) at 3527 cm-1. The presence of a well defined singlet with no evidence of structuration indicates the formation of a unique type of hydroxyl groups on the PC chain during the high temperature dissolution process.

24

.2-

3583

(1)

9 0 DrJ

\ aiz7

.1-

36'00

3 5'00

Wavenumbers

34-'00

(cm-- 1 )

Figure 9. Spectra of PC fracaon (curve ,4) and of pure PC (curve B) in the frequency range 3750-3250 cm "1. The spectra were collected on dilute CH2CI2 soluaons. The formation on the PC fraction of the end-group structures of the type

/0\

/ 0\

OH ,

~

e~f-__~

PC chain --1.6

could account for the spectral features observed in the OH stretching region as well as for the presence of the epoxy ring mode at 916 c m -1. The formation of such end-group structures would imply the presence of hydroxyl-terminated PC chains, formed by chain scission of PC during the high-temperature treatment. The chain scission process of PC can occur either by hydrolysis of carbonate groups (scheme

25 A) or by reaction of carbonate groups with alcoholic functionalities present as impurities in the epoxy resin (scheme B).

SCHEME A

O chain scission, decarboxylation

----.---~ ~

H

+I ~ ~ V v ~ ' L II

+ CO2

II

SCHEME B

ci-I~ T" "J~"~

OH OH I

I

o II

~

CH3 T

O /\ .O-CH2-CH-CH2

/ O\ O-CH2-CH-CH2

H3 -

1.6

-1.6

l

e aO,ox

on

26 /\o C1"!2

/',o

OH ,

~

e~~/F-~

I~-~CH2

+ CO2

-

1.6

In either case, the OH-terminated PC chains (II) can further react with epoxy groups forming epoxy terminated PC chains O /\

O-C~-CH-C~ |+

0 I\

o I\

-CH-CH2

O-C~-CH-C~

H2 -

1.6

To obtain further support for the proposed reaction scheme we have performed quantitative FTIR analysis of the groups involved in the reaction (PC carbonyls and epoxy groups). In particular the epoxy functionalities were determined on chloroform solutions of the 64/36 epoxy/PC mixture using the epoxy ring mode at 916 cmq as an analytical band. To eliminate the interference of the solvent absorptions in the region of interest, spectral subtraction of the solvent was performed. It was found that, at~er the high temperature dissolution process, the content of epoxy groups in the mixture decreased by 7.0 mol % with respect to the starting value. It seems noteworthy that when the pure resin was subjected to the same themml treatment, no reduction in the epoxy group content was detected. An analogous approach was used to detenmne the carbonyl group content in the mixture. In this case the solvent used was CH2CI 2. After the thermal treatment, we found a carbonyl group reduction of 7.8 mol %. Both the epoxy and carbonyl group reductions are consistent with the proposed reaction schemes. From the above analysis it was found that for 100 g of reactive mixture 1.5 x 10.2 mol of carbonyl groups was consumed. On the other hand for the same amount of reactive mixture (100 g), the decrease of epoxy groups was found to be 3.2 x 10.2 mol.

27 The fact that the number of epoxy groups consumed is about twice the number of carbonyl groups seems to indicate that PC chain-seission occurs preferentially through the reaction scheme A. In any ease the formation of structures of type I is extremely important in the subsequent curing processes; in fact the epoxy functionalities at the end of the PC chains will take part in the crosslinking reactions, thus incorporating PC backbones within the epoxy network. After the high-temperature dissolution of PC in the epoxy matrix, the hardener MNA and the accelerator BDMA, were added in the desired proportions. A clear, homogeneous mixture was obtained.

.j o

1.. o m

0 - ~ 4000

-,,

SO'O0 - -

20"00

W a v e n u , u ~ b e r s (err'-- 1)

,600

Figure 10. FTIR transmission spectrum of epoxy~ardener mixture at room temperature prior to curing.

The curing process was carried out at 100~ in an environmental chamber directly mounted in the FTIR spectrometer to monitor the progress of the reaction in real time. The pure epoxy resin was subjected to the same thermal treatment as the blend, prior to the curing process.

28 In Fig. 10 is reported the FTIR transmission spectrum of the epoxy/MNA/ BDMA mixture at the beginning of the curing reaction. The MNA gives rise to several absorptions in the C-H stretching region which further complicate such a spectral range. A well resolved doublet at 1857 and 1781 cm-l appears, due to the symmetric and asynmaetric stretching vibrations of the anhydride carbonyls. Additional strong peaks due to MNA are detected at 1228 cm-l (Vc.o) and at 1083, 943, 929, 916, 899 cm-l (anhydride ring modes). Finally, at 798 crn-l a =CH out-ofplane deformation is found. [35-37]. During the curing process we observed the gradual decrease of the MNA absorptions and the concurrent intensity increase of the peaks arising from ester functionalities (see Fig. 11).

40'00

30'00

20'00

10'00

WAVENUMBERS (era-- 1)

Figure 11. FTIR transmission spectra o f the epoxy~hardener mixture at various curing times.

29 These peaks are located at 2963 cm -1 (v~,,,), 2863 cm -I (E.a,,), 1740

c m "1

(Vc_o, ester), 1454 cm -l (8c~,) , 1398 cm -1 (wcx,) 1267 cm -1 (Vc_o + Vc_c) and 1178, 1155, 1127 cm -l (vr

[36,37].

To follow the conversion of the anhydride groups we monitored the intensity decrease of the well-resolved absorption at 1857 cm "1. From the spectral data it is readily possible to calculate the fractional conversion, ct, of the anhydride groups. The tz versus time curves for the pure epoxy resin (curve A) and for the B 15 blend (curve B) are reported in Fig. 12.

12

9

/

o a

0.6

'

A

I

.....

I

~o -6o"oo o - o - o o o

-iy /

.~a~ 6" o

p

s 0

0

0

-

0.8

O.O'r 0

'

J .....

100

I

200

. . . .

800

time (mirO

Figure 12. Anhydride conversion versus time for pure epoxy (curve .4) and for the B15 blend (curve B). The continuous lines represent the zero order fit; the dotted lines represent the first order fit. Both curves exhibit analogous behaviour with an initiallineartrend followed by a plateau region. Moreover, for pure epoxy a higher initialslope as well as a higher value of the finalconversion arc observed in comparison with the blend.

30 Generally for thermosetting systems the overall reaction rate is a function of the temperature, concentration of reactants, reaction mechanism, and the local microviscosity [39]"

dtz

dt

= A e x p ( - E A/ R T ) f ( a ) f (

1"1,)

(1)

where A is the kinetic Arrhenius factor, E A is the activation energy, R is the molar gas constant, T is the temperature (K), f(ot) is a function of the reaction mechanism and the extent of conversion, and f(rlL) is a function of the local viscosity. Under isothermal conditions the reaction proceeds normally until the molecular weight increases to the extent that the glass transition approaches the cure temperature; f(rlL) becomes important only when the material is about to vitrify. In the absence of diffusion control, the general kinetic equation describing the process is:

do:

dt

where K=A exp

9= A e x p ( - E A/ R T ) f ( a ) f ( r l L )

(2)

(-EA/RT).The simplest expression for ritz) is f ( a ) = ( 1 - a)"

(3)

where n is the order of reaction. Integrating eq. (2) for n=0, n= 1 and n=2, we obtain, respectively:

a = Kt

(4)

- I n ( l - a) = Kt

(5)

6t

1-or

= Kt

(6)

31 Equation (4) fits the experimental data up to a conversion of 0.55 for the pure epoxy and up to 0.60 for the epoxy/PC blend (see Fig. 12). The analysis of the kinetic data according to eq. 5 is shown in Fig. 13. A good linear correlation between-log(l-c0 and t is observed for both the systems investigated. The slopes of the straight lines in Fig. 13 give the first order kinetic constants (K = 0.032 min-1 for the pure epoxy; K = 0.011 for the epoxy/PC blend) from which it is possible to calculate the relative conversion-time curves.

I

I

100

150

"

I v

0

0

0

60

200

time (rain)

Figure 13. -LogO-a) versus time for pure epoxy (curve A) and for the B15 blend (curve B). These curves are shown as dotted lines in Fig. 12. R is apparent that the first order kinetic expression fits the experimental data up to about full conversion for the pure epoxy system and up to 0.8 conversion for the epoxy/PC blend. The analysis

32 performed according to eq. (6) (second-order kinetics) did not yield the expected correlation. From the results of the kinetic analysis some conclusions can be drawn: 1. The data seem to indicate that the presence of PC in the system does not alter the overall reaction mechanism of the curing process. 2. Lower values of the calculated rate constants for both the consumption of anhydride groups and formation of ester groups were found in the epoxy/PC blend in comparison with the pure epoxy system.This marked effect could be ascribed to an increase in the bulk viscosity of the system due to the presence of the dissolved PC.Furthermore the PC may participate to the cross-linking process through its endgroup structures, thus perturbing the kinetics of the epoxy matrix as well as its overall molecular structure. 3. The final conversion of anhydride groups in the epoxy/PC blend is considerably lower than that obtained in the pure epoxy system. Moreover in the latter case a first-order rate equation describes the process over the whole conversion range, thus indicating that the curing reaction never becomes diffusion controlled. Conversely, for the blend, an increasing departure of the theoretical curve fxom the experimental data is observed starting from tx = 0.80. This in turn indicates that at tiffs conversion value the Tg of the system has reached the reaction temperature and the process becomes diffusion controlled. This effect could be ascribed to an increase in the Tg of the blend at any conversion with respect to the Tg of the epoxy matrix because of the presence of the PC component. Such a Tg increase as been demonstrated at zero conversion by DSC measurements.

2.2 Mechanical and Fracture Analysis Dynamic mechanical data for the unmodified resin, compression-molded polycarbonate, and a blend containing 20 % w/w of PC, are compared in Fig. 14. All the materials exhibit a primary tan5 relaxation peak corresponding to the glass transition flg). The Tg of the neat epoxy resin occurs at 170~ while the tan6 peak of PC is detected at 150~

The tan5 relaxation peak of the thermoplastic

component is considerably sharper then that of the epoxy resin. The peak temperature

33

in the blend coincides with that observed hi the neat resin; the peak shape remains highly symmetrical but is considerably broader. Although the Tg'S of the two blend components are quite close, the above observation seem to indicate that no phase separation has occurred during the curing process. In fact, in a phase separated system a less symmetrical band shape of the tan5 relaxation peak would be observed, due to the presence of a second-component peak at lower temperature. Similar results were obtained on the other compositions investigated. SkO

11~

0.6

eC

tO[

~ll E' ( P*I

T.n+

U

11.6

LO 8.2

i

Z.O

t0:"

A

7.0

?.8

74

1

7.0

lOO

so

log ['~

150

P, ~

zoo

1 . . o - o,. -" ,

zso

tZ

9o

to

0.4

lOO

1so

zoo

z1so

Tiin +~

c

i

7.O

!

1,

SO

Figure 14. Dynamic-mechanical spectra at 10 Hz for (A) neat resin; (B) pure polycarbonate; and (C) B20 blend. The fracture behavior of both the unmodified and the PC modified epoxy resins has been examined at low and high strain-rates. At low strain rate (2.5 x 10.3 s -l) two basic types of load-displacement curves are recorded, which correspond to two

34 different types of crack-growth behavior. The curve shown in Fig. 15A is observed for the plain epoxy resin and for blends containing less than 10% w/w of PC (blends B4 and B8). ''

1

I

30 load (N)

j

20

I0

displacement (mm) 0

0.2

I

I

0.4

0.6

Figure 15. Load versus displacement curves for plain epoxy resin (curve A) and BIO blend (curve B).

The load rises linearly with strain up to a maximum value where the crack propagates instantaneously causing a rapid drop in the load. The corresponding fracture surfaces exhibit little evidence of plastic deformation, as will be discussed in detail later. The load-displacement curve shown in Fig. 15B is representative of blends containing 10% w/w or more of PC (B 10, B 12, B 15, B20). Here the crack propagates intermittently in a stick-slip fashion. The load increases linearly up to a critical value and than the crack propagates in a stable manner until the stored elastic energy in the sample decreases to such an extent that crack arrest is allowed. Upon reloading the sample the process of crack-growth is iterated up to the complete failure of the sample. Examination of the corresponding fracture surfaces shows clear evidence of ductility.

35 The critical stress intensity factor K c, is determined from the load-displacement curves according to the equation:

K = crY~f~

(7)

where a is the nominal stress at the onset of crack propagation, a is the initial crack length, and Y is a calibration factor depending on the specimen geometry. For three-point bending specimens, Y is given by Brown and Srawely [40].

I

1.6

1,2

P)

I

I

-

E

Z

0,8

-

0,4

-

w

u

0

I

0

5

"

I

'

10

Blend Composition

I

15

20

PC)

Figure 16. Critical stress intensity factor, Kr at low strain rate as a function o f blend composition.

The values of K c are reported in Fig. 16 as a function of the mount of PC in the blend. It is noted that there is a significant increase of K c with increasing PC content;

36 an essentially linear correlation is found between K c and the amount of the modifier in the blend. The toughening effect of PC is more dramatically evidenced in Fig. 17 where G e data are reported. The G c values were calculated from the values of K c and of the elastic flexural modulus E according to the Irvin relation [41] for linear elastic fracture mechanics:

G =~

(8)

E

In terms of G c, the addition of 20% w/w of PC raises the toughness of the epoxy matrix by a factor of about 7.

600

500

-

400

-

E 300

-

U

200 -

/

J

100 -

0 0

I

I

I

5

10

15

Blend Composition (g) Figure 17. Critical strain energy release rate, G c at low strain rate, as a function o f blend composition.

20

37 In attempting to toughen a brittle polymer, the main goal is to increase the toughness without significantly compromising other important properties such as the elastic modulus. The E values reported in Fig. 18 as a ftmetion of blend composition clearly demonstrate that the improved toughness is achieved without sacrificing the stit~ess of the epoxy matrix. In fact, over the composition range investigated, the modulus shows a gradual but very limited decrease from 3.0 GPa in the neat resin to 2.8 GPa in the B20 blend.

4.5 4~

3.5

13.. r

..

.

I

1

......

I

.........

I

-

-

3.0 -~

w

uJ

2.5

-

2.0

-

1.5

-

1.0

i

0

5

. . . . . . . . . . .

I'

"

10

Blend Composition

I'

15

20

PC)

Figure 18. Elastic flexural modulus, E, as a function of blend composition. Fracture measurements were also carried out under impact conditions in order to evaluate the toughness of these materials under rapid loading. The Kc and G e values are reported as a function of blend composition in Figs. 19 and 20, respectively. In

38 this case the K e values were obtained as previously, using eq. 7, while the G e values were estimated by energy measurements according to the following equation:

U

(9)

G = BW@

where U is the fracture energy corrected for the kinetic energy contribution, B and W are the thickness and the width of the specimen respectively and ~ is a calibration factor which depends on the length of the notch and the size of the sample. Values of 9 were taken from Plati and Williams [42]. Apart from a decrease in the absolute values of Ke and G c the general behavior of the impact toughness parameters is analogous to that observed in the low-speed tests.

1.5

I

I

I

5

10

15

1.0-

0.5

0

20

Blend Composition (% PC) Figure ] 9. Critical stress intensity factor, K c, under impact conditions as a function of blend composition.

39 350

300

-

250

-

200

-

150

-

100

-

(M

E

--.j w

o

(.t

500i~ 0

/

I

I

I

I

5

10

15

20

Blend Composition (g)

Figure 20. Critical strain energy release rate, G• under impact conditions as a function o f blend composition.

In particular Kr and Gr increase by factors of about 2 and 5, respectively, compared with the values for the neat resin. A decrease in toughness on increasing the loading rate is a general phenomenon related to the reduced capacity for viscoelastic and plastic deformation which polymeric materials exhibit when the strain rate is enhaneexi. However, in our case it is interesting to note that, even if the strain rate is increased by about 5 orders of magnitude, the observed decrease in the toughness parameters is relatively modest (e.g., 40 % for the B20 composition). The considerable increase in toughness found in the PC-modified epoxy resins can be ascribed to an improved capability of localized plastic deformation of the epoxy/PC network. This enhanced capacity of deformation is due to the particular kind of molecular structure of the network developed during the curing process.

40 As previously mentioned during the high-temperature dissolution process, the PC component interacts chemically with the uncured epoxy resin with formation of epoxy end group. Such groups located at the ends of PC chains can take part in the curing reaction thus incorporating PC chains into the epoxy network. The resulting molecular structure can be sketched as:

( dissolution 220uC

~ 5hr=

curing postcuring

_

The PC molecular segments are longer and more flexible than the segments of a simple epoxy network and can be more easily deformed under loading. The fracture toughness data have been interpreted in terms of the morphological analysis of the fracture surfaces obtained on samples tested at low and high strain rates. It has been found in earlier investigations [43-46] that several characteristic features can be observed on fracture surfaces of epoxy resins, especially when stickslip propagation takes place. These features may fall into three main categories: an initiation region followed by crack arrest lines, a region of slow crack growth, and an area of rapid crack growth which covers the remaining surface of the sample. The scanning electron micrographs of the fracture surfaces of samples tested at low strain rate (Fig. 21) illustrate the above features. In particular, we note that, prior to the crack arrest line AB, all the samples exhibit a smooth and relatively featureless surface which can be associated with fast crack propagation,. Beyond the AB line, the pure epoxy (Fig. 21A) as well as the

41 blends B5 and BS, for which no stick-slip behavior was observed (micrographs not shown), display only the presence of fine markings extending from a restricted area.

Figure 21. SEM micrographs offracture surfaces obtained at low strain rate." A) pure epoxy resin; B) BIO blend; C) B15 blend; D) B20 blend. At high magnification (Fig. 22) these markings appear as wave crests and arise from adjacent sections of the crack front following paths at slightly different levels. On the other hand, when the crack propagation occurs by a stick-slip process, after the crack arrest line a well defined slow-growth region (area ABCD in Figs. 2 I B, 21C and 21D) is observed. The size of this area increases markedly with increasing PC content in the blend. A closer examination of this region at higher magnification

42 (see Fig. 22) reveals the occurrence of V-shaped features that result from events occurring during the arrest and re initiation of crack propagation.

Figure 22. High magnification SEM micrographs offracture surfaces obtained at low strain rate: .4)pure epoxy resin; B) B20 blend. Following the slow growth region, there is a transition to an unstable initiation region where the crack accelerates, giving rise to the formation of fine longitudinal lines, approximately parallel to the crack direction. Note that the crack direction is from left to right in the mierographs. The above observation account for the fracture toughness results and, at the same time, give a clear picture of the events occurring in a stick-slip process: after crack arrest a localized plastic zone develops at the crack tip upon loading, and the crack becomes blunted. This blunted crack grows slowly until sufficient strain energy is stored, through continued loading, to force the crack to propagate rapidly through file undeformed material.

43 Thus the slow growth region observed in figure 2 l b-d, defines the size of the plastic zone at the crack tip. Such a plastic zone is most likely generated via a shearyielding rather than a crazing mechanism. Indeed several authors [7, 9, 47] have concluded that crazing does not usually occur in epoxy materials. Therefore, it might be assumed that localized yielding at the crack tip with consequent notch blunting is the major source of energy dissipation during fracture in both the unmodified and the PC-modified epoxy resins. This mechanism is far more active in the B 10, B 15, and B20 blends because of the rather extensive presence of PC chains within the epoxy network. In fact as shown in Figs. 2 l b-d, an increase in size of the plastic zone is observed on increasing the PC content in the blends. Fig. 23 shows SEM mierographs of some of the investigated samples fractured under impact conditions.

Figure 23. SEM micrographs offracture surfaces obtained under impact conditions: A) pure epoxy resin; B) BIO blend; C) B20 blend.

44 Examination of these surfaces at low magnification did not reveal features characteristic of a stick-slip process because, at high strain rate, this type of crack propagation mechanism is completely suppressed. It is clear that, even at very high magnification (2500 x), no evidence of a dispersed second phase could be detected in the blends, thus confirming the dynamic-mechanical results which, for a B20 composition, indicated the occurrence of a single-phase, homogeneous system. Attempts to selectively etch the fracture surfaces with a solvent for PC (CH2CI 2) failed to reveal any feature distinctive of PC removal. The fracture data were used to calculate the degree of crack tip blunting occurring when these materials are fractured at low strain rate. In order to estimate such a parameter the stress distribution around a blunt crack has to be taken into account. It can be shown [48] that, for a crack under an applied stress of cro, the stress normal to the axis of the crack at a short distance r ahead to the crack tip is given by:

or= Cro~-~r (1 +l + p /p2/rr) "

(lO)

where p is the crack tip radius and a is the crack length. Assuming that the fracture occurs when a critical stress r is reached at a distance r = c, eq.10 becomes [48,49]:

c r ,~~ = (l+p~/2c) "~ cr~/2av (1 + p, / c)

(11)

The term cr 2 ~ ~ can be considered as the critical stress intensity factor Kle for a "sharp" crack, and cr~/no as the stress intensity factor K~ for a blunt crack. Hence eq. 11 may be rewritten as:

K~ = ( l + p , / 2 c ) ~'~ Kk (l+p/c)

(12)

45 This equation relates Kr to the radius of a blunt crack Pc; its validity may be checked by measuring the variation of K~ with pC. Direct evaluation of Pc is very difficult, especially for thermosetting systems. However, as reported by Kinloch and Williams [48], this problem may be circumvented assuming Pc to be equivalent to the crack opening displacement 8c:

(13)

where ay,t is the tensile yield stress and Sy is the yield strain. The above equation has been used to calculate 8~ or p~, and hence the degree of crack blunting. The resulting p~ values are reported in Fig. 23 as a function of blend composition. 15

I

I

,,

I

10 E U

o

I

0

5

I

10

Blend Composition

I

15

20

(~ PC)

Figure 24. Crack-np blunting, Pc, as a function of blend composition.

46 It is noted that the degree of crack tip blunting increases sharply starting from a PC content of 10% wt/wt. This behaviour further confirms that blunting takes place only when stick-slip propagation occurs. In terms of yield behavior, the appearance of such a type of crack-growth mechanism might be viewed as a consequence of the reduction of the yield stress, whiell enlmaces the ability of the material to plastically deform in the vicinity of the crack tip with a consequent increase in the fracture toughness parameters.

3. TGDDM/PEI system For this particular blend system the thermoplastic component was dissolved in a common solvent. In particular the PEI was dissolved in CH2CI2 and mixed with the epoxy resin at room temperature. After complete dissolution of the TGDDM, the solvent was distilled off at 60~ vacuum at 100~

The last traces of CH2CI2 were eliminated under

obtaining a clear, viscous solution. The DDS hardener was added

at 120~ under vigorous mechanical stimng up to complete dissolution. The mixture was poured in an open steel mould and degassed in a vacuum oven at 100~ for 5 h. Finally, the blend was cured for 16 h at 120~

2 h at 150~

2 h at 180~ and

postcured for 4 h at 200~ For the blend containing 30 phr of PEI, due to the very high viscosity of the resulting mixture all the components were dissolved in CH2C12; solvent removal and degassing were performed simultaneously in the vacuum oven at 100~

Blend

compositions and codes are reported in Table 3. DSC measurements were performed on several of these mixtures of varying compositions and are reported in Fig. 25. As in the case of the epoxy/PC system, a single glass transition temperature (Tg) is observed in all the cases, which increases by increasing the PEI content in the mixture. The Tg values are shown in Fig. 26 as a function of composition. The above results indicate that, in the composition range investigated, the PEI is molecularly dispersed in the TGDDM forming a single-phase homogeneous system. It

47 is worth noting that, when the DDS hardener is dissolved in these mixtures, they still remain homogeneous and visually transparent. Table 3.

Codes and composiaons of the investigated TGDDM/PEI mixtures. Code

TGDDM

DDS

PEI

PEI

A0

76.9

23.1

-

-

A5

74.1

22.2

3.7

5.0

AI0

71.5

21.4

7.1

10.0

A15

69.0

20.7

10.3

15.0

A20

67.0

20.0

13.3

20.0

A30

62.6

18.7

18.7

30.0

(%)

(%)

,,,

(%)

oT X Iii

-50

-25

0

25

Temperature

!

I

I

5O

75

100

('C)

Figure 25. DSC thermograms in the temperature range between-50~ and lO0~ for,4) AO; B) AIO; C) A20 and D) A30 blends.

48 I

320

310

I

I

-

300 -

290

-

280

-

270

-

260

'

I--

0

I

.

10

20

I' 30

40

Blend Composition [g PEI) Figure 26. Glass transition temperatures, Tg of the uncured TGDDM/PEI mixtures as a function of composition. The residual heat of curing, AI-Ir normalized to the epoxy resin content, is reported in Fig. 27A as a function of blend composition. We observe a linear increase of AI~ with increasing the PEI content in the blend. After the postcuring step. the neat resin shows no residual M-Ir. Conversely, for all the blends, AHr is considerably lower than that observed in the cured samples but still well detectable (Fig. 27B). Moreover a linear trend, approximately parallel to fl~at observed for the cured samples, is found between AI~ and P EI content. These results can be interpreted considering that, as it will be shown later, the PEI phase separates during the early stages of the curing process. It is likely fllat this PEI phase incorporates a small amount of unrcacted TGDDM which can act as plasticizer [50,51].

49 I

200

150

Ik,,,

!

I

-

100

"r'

._

125 The PEI was dissolved in CHCI 3 and BMI and DDS were added to the solution under vigorous mechanical stirring. The solvent was removed under vacuum and the mixture was placed in an open steel mold and cured at 160~ for 7 h. Postcuring was performed for 2 h at 180~ and 4 h at 200~

The codes and

composition of the investigated blends are reported in Tab. 1.

Table 1. Codes and compositions of the investigated BML'PEI mixtures.

Code

BMI

DDS

PEI

PEI

(wt %)

(wt % )

(wt %)

(p.h.r.) a

A0

76.9

23.1

-

A10

71.5

21.4

7.1

10.0

A20

66.7

20.0

13.3

20.0

A30

62.6

18.7

18.7

30.0

.

.

.

.

aparts for 100 parts of BMI.

The fracture behavior of the net BMI resin and of the PEI modified resins was examined under impact conditions using a Charpy instrumented pendulum. The data were analyzed according to the linear elastic fracture mechanics approach. The parameters Ko (critical stress intensity factor) and G~ (critical strain energy release rate) were calculated by means of the equations: =

rrqS

l)

where cy is the nominal stress at the onset of the crack propagation, a is the initial crack length and Y is a calibration factor depending on the specimen geometry. U G =~~BWr

2)

126 where U is the fracture energy corrected for the kinetic energy contribution, B and W are the thickness and the width of the specimen respectively, and 9 is a calibration factor which depends on the length of the notch and size of the sample. The values of 9 were taken from Plati and Williams [11]. Figs. 1 and 2 show the values of K c and G c as a function of blend composition. Both the parameters increase significantly with increasing the amount of PEI in the blend; in particular G c increases three times with respect to the reference BMI matrix. Moreover no reduction of the elastic modulus of the material was observed for all the investigated compositions.

t

0.8

...........

I .........

.

-

f

N

E Z

1.

0.6

-

z; (

,,r 0.4

0.2

. . . . . . . . . . .

0

I .

5

.

.

.

.

.

[ . . . . . .

10

I

i,

15

Composition (PEI Wt Figure 1. The critical stress intensity factor, K c, determined under impact conditionsjbr the BMI/PtH system as ajbnction of blend composition.

i

2O

27 250

. . . . . .

I

. . . . .

I

. . . .

1 .....

200 -

E

150

-

""3 u

100

50

0

r

0

'

"

i

5

. . . . . . . . . .

i

10

.

.

.

.

.

i

15

20

Composition (PEI W t %) Figure 2. The critical strain energy re~ease rate, G c, determined under impact conditions for the BMI/PEI system as a function o f blend composition.

The morphological analysis performed on the fractured surfaces clearly indicates that the system is completely phase-separated [12]. However the f'mal morphology strongly depends on blend composition. At low PEI content (see Fig. 3) small PEI domains (about 2 l.tm), uniformly distributed within the BMI matrix cohexist with very large domains whose size ranges from 50 ~ n to 100 ktm. Upon etching these surfaces with boiling vapours of CH2C12, the small domains disappear while the larger ones are not removed by the solvent but become more porous. It is likely that in these regions a phase inversion has occurred. By increasing the PEI content in the blend the larger domains tend to

128 increase in size and to coalesce, and a very limited number of smaller domains of pure PEI is observed (see Fig. 4).

Figure 3. SEM micrographs of the fracture surfaces of samples broken under impact conditions Jbr A) unetched A 10 blend; B) etched A 10 blend.

129

Figure 4. SEM micrographs of thefracture surfaces of samples broken under impact conditionsfor A) unetchedA20 blend; B) etchedA20 blend. Finally, at a PEI concentration of 30 phr, the morphology of the whole fracture surface closely resembles that of the large domains observed at lower PEI content (see Fig. 5).

130

Figure 5. SEM micrographs of thefracture surfaces of samples broken under impact conditionsfor A) unetchedA30 blend; B) etchedA30 blend. The above morphological analysis gives an insight into the fracture behaviour of this blend system. In the regions where PEI is the dispersed phase, the fracture occurs by brittle failure of the BMI matrix with the PEI domains bridging the crack and delaying its

131 propagation. Conversely, in the regions where PEI forms the continuous phase the failure may occur within the thermoplastic phase and the crack propagates around rather than through the BMI domains. Therefore, yielding of the thermoplastic continuous phase is the main toughening mechanism. These two deformation mechanisms operate simultaneously in the BMI/PEI system; their contribution to the over-aU fracture tougheness depends strongly upon the blend composition.

2. BMI thoughened by reactive liquid rubbers It is generally accepted that the addition of a reactive liquid rubber can improve substantially the tougheness of thermosets like epoxy resins. Only a few papers on the modification of BMI resins with rubbers have been reported so far, whereas rubber-modified epoxies have received considerable attention in the literature [ 13, 14]. St. Clair and St. Clair have reported an increase in the fracture tougheness of nadic-terminated polyimides by addition of amine-terminated silicone rubbers [15], whereas Varma et al. [16] have reported an increase in shear strength of BMI resins by the addition of amine-terminate~i rubbers. Shaw and Kinloch observed an improvement in the fracture energy of BMI resins by modification with carboxylterminated butadiene-acrylonitile (CTBN) rubbers [17]. Unlike with epoxy resins, CTBN rubbers are not compatible with BMIs even at high temperatures (110~ However, it has been claimed that during curing between 170~ and 210~ the CTBN rubber reacts via the backbone double bonds with the BMI resin, although no experimental evidence has been provided to conf'mn such a hypothesis. In any event, at the end of the curing process a microphase separated structure similar to CTBN-modified epoxies, is observed. The main disadvantage of this approach is the marked decrease in the high temperature mechanical properties. The elastic modulus at 250~ shows only 0.21 GPa for the 50/50 BMI/CTBN rubber system, which also indicates a significant reduction of the glass transition temperature. As a result of the oxidative sensitivity of the CTBN rubber, the oxidative stability of

132 the BMI/CTBN system is low. Nevertheless, the fracture tougheness enhancement makes this system interesting for adhesive applications. Recently new bismaleimides containing ether linkages were prepared and characterized [18]. Mixtures of two of these products form an eutectic mixture and become easy to process owing to their low melting point and long pot life. Takeda and Kakiuehi [18] studied the modification of these mixture systems with CTBN rubbers. In particular they investigated the mechanical and thermal properties as a function of the amount of and the composition (aerylonitrile content) of the added rubber. The chemical formulas of the investigated BMI resins are reported below: o

o

6

6

where R is:

4,4'-bismaleimidodiphenylmethane(BDM)

2,2-Bis [4-(4-maleimidophenoxy)phenyl]propane (BPPP)

Bis [4-(3-maleimidophenoxy)phenyl]sulfone(3,3'BPPS)

133 In Fig. 6 are reported the results of flexural tests of a resin modified with CTBN containing 17 % of AN. The base resin was a 50/50 wt/wt mixture of BDM and 3,3'-BPPS. The flexural strength increases with increasing the amount of CTBN up to 50 phr and then levels off. However a marked reduction of the flexural modulus is observed by increasing the CTBN concentration.

12

.

.

.

.

I. . . .

l

i

,,,

I

500 "3'1 --,--.,,

E E

(11) x

c,.

,

- 400"

-

0

0.

- 300

,m, I,,.., 4.,.I

m

- 200

I,.,

x

=_.

r or)

""

3 3

I.I_ !

I

o

20

I

40

CTBN

'

'!

. . . . . .

60

I

. . . .

80

Concentration

I

100

100

'

120

(phr)

Figure 6. Flexural strength and flexural modulus as a function of CTBN concentration. After Takeda and KakiuchL

The fracture energy (Gc) was found to increase with the rubber content while decreased by enhancing the AN content in the rubber (see Figs. 7 and 8 respectively). The AN content had a strong effect on the solubility of the rubber in the uncured resin and, as a consequence, on the final morphology of the cured

134 material. In the absence or at low AN coments (up to ~ 20 % wt/wt) the rubber is insoluble in the resin and the blend shows a microphase separation. At higher AN values, the CTBN becomes soluble and no phase separation occurs upon curing. In these conditions the rubber modifier is far less effective in inducing energy dissipation mechanisms [18]. The significant improvement in tougheness observed for phase separated systems, was interpreted as a consequence of a good interfacial adhesion among the phases. In turn such an effect was ascribed to a chemical interaction occurring between CTBN and the resin matrix.

350

t

!

~,

I

....

I

.........

I

300 250 0r

E

(5

200

-

150 -

7

100 50-

'

0

'

'

i ...................

20

! . . . . . .

40

I

60

. . . . . .

I'

80

'

!

100

120

CTBN Concentration (phr)

Figure 7. Fracture Energy (Go) as a function of rubber concentration. After Takeda and Kakiuchi.

135 The authors [18] determined the concentration of carboxyl end-groups of CTBN prior and after a pemixing step, and they found that such a concentration hardly changed. Therefore, in analogy to what happens for polyisoprene radically crosslinked in the presence of catalytic amounts of bismaleimide, they suggested that the crosslinked network contains linkages between the allylic carbon of the rubber and the maleimide double bonds [19].

1 6 0

'

140

-

,,

,

I

,

I

,

.

l

.,,I

I

.....

E 120 -

0 100

80

.... , 0

'

5

I

10

.....

I

. . . . . .

15

I

.

.

.

20

.

.

I

25

"

'~

30

AN Content (wt ~1 Figure 8. Fracture Energy (G~ as a function of AN content in the rubber. After Takeda and KakiuchL

The authors of the present chapter recently reported [20] on the modification of a bismaleimide resin commercially available as Kerimid FE 70026 (RhonePoulenc) which is a mixture af BMI and tolylene-BMI in the weight ratio 60/40.

136

0

Tolylene-BMl

This resin is partially oligomerized by a diamine to improve its processability. In fact it is a glassy solid at ambient temperature and becomes fluid at about 50~

no free primary amine groups are present in the formulation. The

exact molecular structure of this product is proprietary and this limits to some extent the spectroscopic characterization. Nevertheless Kerimid 70026 has been used in place of other, better characterized BMI monomers, owing to the fact the the rubber modifier can be readly dissolved in such a matrix. The toughening agent was a maleimido-terminated butadiene-acrylonitrile copolymer (ITBN) whose preparation and spectroscopic characterization is reported in the second chapter of this book. Initially a kinetic analysis as a function of temperature was performed on the pure polyimide as well as on a typical blend composition (85/15 wt/wt) in order to investigate the effect of the rubber modifier on the curing kinetics of the resin and to identify possible chemical interactions among the blend components. Real-time FTIR spectroscopy measurements were carried out in the transmission mode, placing a thin film (1 - 5 ~m) of the product in a temperature chamber mounted in the spectrometer. In Fig. 9 are reported the the spectra of the neat Kerimid collected at 160~ in the frequency range 3250 - 2700 cm -1. Two peaks characteristic of the maleimide double bonds are located at 3167 cm -1 and 3100 cm -1. The former has been attributed to the first overtone of the C=C stretching vibration, while the latter is due to the C-H stretching mode of the bismaleimide unsaturation [21-23]. Both the peaks are found to decrease with reaction time, thus confirming the gradual disappearence of maleimidic double bonds. However the 3167 cm -1 peak

137 is too low in intensity to afford reliable quantitative evaluation. The 3100 cm -1 peak has a reasonable intensity but suffers from extensive overlapping with the peak at 3070 (aromatic vC_H). The spectral data were thus analyzed by means of subtraction spectroscopy, whereby the spectrum at time zero is subtracted from those collected at longer times:

/'/']

lD

o

C O JE~

29.31

217

.1 -

3100

0 It) _13

30,37

.05

2'oo

30'00

Wavenumbers

(cm--1)

28'00

Figure 9. Real time spectroscopic monitoring of the curing reaction of the neat Kerimid resin in the frequency range 3300- 2700 ctrr'. Spectra collected at 160~

4 = 4-SF.Ao

138 where the subscripts s, t and 0 denote the absorbance of the subtraction spectrum, of the spectrum collected at time t and of the specmun collected at time zero, respectively. The subtraction factor, SF, allows to compensate for differences in thickness between the spectra collected at times zero and t. Its value is obtained by reducing to the baseline an internal thickness peak, i.e. a peak which is invariant with the reactants' conversion (in the present case the aromatic absorption at 1514 cm-1). It has been found that the thickness variation during the process is very limited, so that SF is always close to unity.

.01

-

_

tO

i

i_ 0

--.01

-

I .......

3 0o 1 .

3200

.

.

.

.

.

.

.

.

.

2!72

2965 .

.

.

.

3000 i

Wovenumbers

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

Ccm-1)

.

.

.

.

.

2800 i

Figure 10. Spectral subtraction analysis of the data reported in Fig. 9.

139 The main advantage in using spectral subtraction in the analysis of kinetic data rely on the fact that peaks not affected during the process (the interfering absorption at 3070 cm -1 in the present case) are compensated for and hence completely removed in the subtraction specmma. Conversely positive absorbances from the zero base line generally reflect molecular structures that are formed during the process, while negative absorbances reflect structures that are lost. The subtraction spectra reported in Fig. 10 in the frequency range 3200 2700 cm -l clearly show the gradual development of a completely resolved negative peak of BMI at 3100 cm -], for which a linear and consistent baseline can be identified. The absorbance values so evaluated can be directly used to determine the BMI conversion as a function of time.

0.6 !I.....

i

i

I

.............

!

. . . . . . . .

1

_

180~ 0.4

,f

o ~ o , ~ M o-~:~ o

1 0 C0

_..;

6o'c

o

0.2

f

0

~

I

100

. . . . . .

I

200

I. . . . .

300

I

400

. . . . . . . .

I

500

600

Time (min) Figure 11. The conversion o f bismaleimide double bonds as a function o f the reaction #me.for the neat Kerimid resin. Temperatures as indicated.

140 In Fig. 11 are reported the conversion, (x, versus time curves for the neat Kerimid resin at three different temperatures (160~

170~

180~

All the

curves show an initial linear trend whose slope increases by increasing the temperature. At longer times the reaction rates decrease substantially and the curves approach a plateau region. The fmal conversion increases with increasing the temperature. It is noted that at 180~ the final conversion values do not excede 50 %; this is due in part to the rigidity of the resin, whose Tg increases rapidly with conversion. Thus at o~ values close to 0.5 the glass transition reaches the reaction temperature and the curing process is frozen in. In order to achieve a more complete cure it is necessary to process the resin at temperatures exceeding 200oC.

0.8 .!...................... i ................

0.6

I

~

I ..............

180"0

0

~. . . . . . . .

170"C 0

O_ 0

0.4

1

50 ~ 0

0

0.2

~ f O~r0

...............................

i

200

................................

l ....

400

......... i ......................................[' 800 600

Time (min) Figure 12. The conversion o f bL~maleimide double bonds as a function o f the reaction timeJbr the 85/15 Kerimid/ITBN blend. Reaction temperatures as indicated.

141 When the ITBN component is premixed with the Kerimid matrix, noticeable effects are detected in the curing kinetics. The ot-t curves relative to a blend containing 15 % wt/wt of ITBN at 150~

160~

170~ and 180~ are reported

in Fig. 12.

0.6

. . . . .

I

,

.I

. . . . . . . . . ~. . . . . .

t,

,

I

,

,

160~ A

0.2

tl 0

T

0

I

1

100

200

i

300

!

"

:

400

!

'

500

600

Time (min)

0.6

j

........

i 170oC

t

,

t,,

,

i

.

,

i

0.4

-

0.2

-

0 0

100

200

Time (min)

300

400

142 0.8 I| . . . . . .

I .....

l

....

I ....

,

..

J

..

180'C

0.6 l

B

0.4-

0.2-

0

. . . . . . . . 0

,.

50

. . . . . .

i

....

100

,

. . . . .

150

,

200

.......

i~

250

300

Time (min)

Figure 13. Comparison of the bismaleimide conversion profiles relative to the neat Kerimid resin (curves A) and to the Kerimid/ITBN 85,/15 blend (curves B). Reaction temperatures as indicated.

To highlight the differences the kinetic curves relative to neat Kerimid resin (curve A) and to the blend (curve B) at 160~

170~ and 180~ are compared in

Figs. 13. It is readly apparent that both at 160~ and at 170~ the curing process in the blend is considerably retarded; in particular the initial linear trend occurring in the neat resin is almost completely suppressed in the blend. Moreover, for the blend, owing to the slower reaction rate, a plateau region was not reached in the investigated time range although it appears that the ot-t curves of the blend tend to the same final conversion values of the neat resin. However the above retardation effect is found to depend strongly on the reaction temperature: it decreases at 170~ while at 180~ the situation is completely reversed. Here both the reaction rate and the fmal conversion are higher in the blend than in the neat resin. (see Fig.

13).

143 The retardation effect observed at 160~ and 170~ might be accounted for by a viscosity increase occurring when the ITBN rubber is dissolved in the neat resin. This in turn causes a reduction in molecular mobility of the maleimide functionalities involved in the crosslinking process. As the temperature increases such an effect is reduced because of the decreasing viscosity of the blend system and, more importantly, because of the higher rate of production of primary radicalic species which renders the system less sensitive to molecular mobility effects. The observation that at 180~ an opposite effect is found on both the reactio~ rate and the fmal conversion is not easily accounted for. One possible explanation could be the occurrence of finther reaction steps which are inactive at lower temperatures. These processes involve chemical interactions between the rubber unsaturations and the bismaleimide double bonds and have the net effect of complicating the overall reaction mechanism and of accelerating the chain addition process through which the BMI network is built up. Spectroscopic evidence of the occurrence of such reactions at 180~ is discussed below. The problem of the possible chemical interactions between the ITBN rubber and the thermosetting matrix was investigated spectroscopically. In principle these interactions may occur either through the maleimide-end groups of ITBN or by reaction of the double bonds along the ITBN backbone. The first kind of interaction is not amenable to spectroscopic analysis since the rubber end-groups yield the same signals as the maleimide groups of the matrix, and the contribution of the two species cannot be separated. However, since alifatic bismaleimides are more reactive than their aromatic counterparts, it is very likely that the BMI end-groups of the rubber partecipate to the cross-linking process through which the kerimid network is build up. With respect to the unsaturations present along the rubber backbone, it is well known that, in general, they may have three different configurations:

144 ~9 C H 2 \

/H /C--C\ H CH2-~

---CH2\ /CH2---/C=C,, H

H

trans-l,4

cis-l,4

- - ~ C H 2 ~ H ....... CH 11 CH2 1,2-vinyl These configurations yield three distict characteristic frequencies due to the out-of-plane deformation of the =C-H bond at 730, 970 and 915 cm -1, respectively. The ITBN transmission spectrum reported in Fig. 14 shows the presence of two well resolved absorptions at 968 cm -1 and at 915 cm -1 while the characteristically broad peak at 730 cm -1, distinctive of cis-l,4 configurations is absent. Thus in the ITBN copolymer only trans-l,4 and 1,2 vinyl configurations are present, with the former being largely predominant. The knowledge of the ratio gtrans/l~1,2v [24] allows one to evaluate quantitatively the relative population of the above configurations by using the relationship:

Ctrans C],2v

~ ~

Arrans A1,2v

In the present case it is found that 82.3% of monomeric units are present in the trans-l,4 configuration and 17.7% in the 1,2-vinyl configuration. By using spectral subtraction spectroscopy it has been possible to isolate the specmnn of ITBN in the region of interest from that of the blend. This was accomplished by digitally subtracting the contribution of the matrix from the blend spectrum.

145 J

g68

1

g15

.5 j ~ o'o0

3 0 0 0.

. . . . . . .

2 0 i0 0 .

.

.

.

96o

.

.

1'

0~'0 0

Wavenumbers (cm- 1) Figure 14. The ITBN transmission spectrum in the frequency range 4000 450 cm-'. The inset evidences the frequency range where the characteristic group frequencies of the double bonds occur.

The situation is represented in Fig. 15, where curve A refers to the initial blend specman taken at 160~

while curve B represents the result of the spectral

subtraction; curve C is the ITBN specman at ambient temperature and is reported for comparison. It is noted that traces B and C are almost coincident, apart from a sligth broadening of the peaks in the former spectnnn which is due to a temperature effect. The spectral data of Fig. 15 demonstrate the reliability of such an approach but, in order to be able to monitor the fate of the minor component in the blend an appropriate criterion is needed to choose the reference spectrum to subtract at times other than zero. This is because both the blend and the reference spectra change quite substantially with time due to the crosslinking process, and the rate of change of the two spectra are quite different.

146

"1 968

I .,3 915

(.1 to

-s o

~t .2

10'00 950 900 Wavenumbers (cm--1)

Figure 15. ,4) The B15 blend spectrum in the frequency range 1050- 850 cm'; B) The subtraction spectrum Blend- Neat Resin; C) The ITBN spectrum. Thus, in performing the spectral subtraction analysis we chose to use a matrix speetrtun and a blend spectrum having coincident values of conversion; the situation is schematically represented in Fig. 13" to obtain the difference specmma representative of ITBN at time t 1, the matrix spectrum corresponding to the point r I was subtracted from the blend spectrum corresponding to the point s I. In this way "clean" results were obtained over the whole time range investigated. It is noted however that this approach cannot be used for the kinetic curves taken at 180~

for times higher than 50 min. This is because at longer times the

conversion values of the blend excede the maximum conversion of the neat resin

147 and a reliable reference spectrum is no longer available. In these cases the spectral components at 968 and 915 cm" were separated by a curve fitting algorithm; the details of the calculations are reported in Ref. 23. In Fig. 16 are reported the conversions of the trans, 1-4 unsaturations and those of the 1-2, vinyl unsaturations of the ITBN rubber as a function of the reaction time for the process carried out at 160~

and 170~

(Fig. 16 A) and at

180~ (Fig. 16 B). First it is noted that both at 160~ and at 170~ the conversion of the two unsaturated species is close to zero, which indicates that in this temperature range the double bonds along the rubber are not involved in any kind of chemical interaction.

,I

0.8

0,6

-

0,4

-

0.2

L

I

I

I

I

!

I

m

m -

[]

12 0.0

-~

-0.2 0

"H

~

U

I

1

I

50

100

150

,

I 200

,,

, 250

time (min)

9

J 300

I

!

350

400

450

148 0.8

0,6

I

,

I

-

0.4

-0.2

L

I

.//

o~

~

9o

0.2

0.0

,

Oo~176 ~ 4

I 0

~

o

l

.......... 50

, 100

, 150

200

time (rain)

t~gure 16. The conversion o f the trans, 1-4 and o f the vinyl 1-2 unsaturations o f lTBN as a function o f the reaction time. m,["l= 160~ 1700(7..; 0 , 0

-- 180~

A, A =

The open symbols rejkr to the 1-4 trans unsatura-

tions, the solid symbols to the 1,2-vinyl.

A differem picture is found at 180~

Here, aider an induction period lasting

approximately 50 min. the conversion of the above species starts to increase substantially with an approximately linear trend and reaches values of about 60 % by the end of the process. No evidence of a plateau region is found in the investigated time interval: the process through which the rubber unsaturations are consumed would have continued fimher, perhaps to completition, at longer times. Another relevant observation is that the data points relative to both the types of

149 unsaturations can be accomodated on a single curve, which indicates a similar reactivity of both the trans-l,4 and the 1,2-vinyl unsaturations. At this point it is worth to compare the kinetic profile relative to the BMI double bonds (Fig. 13) with those of the ITBN unsaturations (Fig. 16 B). It is found that the rubber double bonds start to react at about 50 min when the BMI conversion is already high (45 %). At this point the BMI reaction rate has already slowed down considerably, while in the region were it was at its maximum the ITBN unsaturations were inactive. Furthermore in the time range where the reaction rate of the rubber double bonds is steady, the BMI reaction rate gradually reduces to zero; towards the end of the process the BMI conversion remains constant while that of the rubber unsaturations continues to increase linearly. In summary it seems that the above processes are not directly correlated, as if they were following two different and independent reacction pathways. A crosslinking process of the ITBN rubber within phase separated rubber domains initiated by catalitic amounts of BMI and/or by the maleimide end-groups of the rubber would explain this effect. In this instance the induction period would represent the time necessary to reach a critical concentration of radicalic species in the rubbery domains. This kind of mechanism has been demonstrated for polyisoprene radically cross-linked in the presence of small amounts of bismaleimide. On the other hand this mechanism would not account for the acceleration of the curing process of the matrix observed in the blend at 180~ and for the increase of the final BMI conversion compared to that in the neat Kerimid resin. In fact chemical processes confined into the rubbe13, domains of a phase separated system would hardly affect the over-all curing mechanism of the BMI continuous phase. Furthermore a morphological analysis of the fracture surfaces of blends of different composition carried out by scanning electron microscopy [20] did not reveal any evidence of a dispersed second phase, even at very high magnification. Thus the electron microscopy analysis indicated the occurrence of a single phase, homogeneous system upon the curing and postcuring processes.

150 The experimental data just discussed cannot be considered conclusive and further investigations on the molecular structure realized upon curing in such a complex network, possibly employing other spectroscopic techniques, are in order to fully account for the experimental observations. However, at temperatures of 180~ and above, extensive chemical interactions between BMI and ITBN can be anticipated which involve both the maleimide end-groups and the backbone unsaturations of the robber. To test the ability of the ITBN rubber in improving the tougheness of the Kerimid matrix, a series of blend compositions were prepared by first dissolving the ITBN rubber into the resin at 120~ for 30 min, and degassing under vacuum for additional 30 min. At the end of this step a clear, visually homogeneous mixture was obtained. The mixture was then poured in a glass mold, cured at 180~ for 5 h and postcured at 220~ for 2 h. The codes and compositions of the investigated blends are reported in Tab. 2

Table 2. Codes and compositions of the investigated blends" Code

BMI

ITBN

(wt%)

(wt %),,,

K0

100

-

K4

96

4

K8

92

8

K10

90

10

KI5

85

15

The fracture properties of such blends were investigated at low (1 mm/min) and high (1 m/sec) rate of deformation. The parameters K e and G~ were calculated using Eqs 1 and 2.

151 In Fig. 17 the K c values are reported as a function of blend composition; for both the testing conditions K e increases linearly with increasing the rubber content up to a maximum of about two times for a 85/15 blend composition. Similar results are found when the fracture tougheness is expressed through the parameter G e (see Fig. 18). In this case, for the low speed tests, an increase of about three times is achieved with respect to the value of the neat resin for the 85/15 blend composition. As shown in Fig. 19 a modest reduction of the elastic modulus is brought about by the addition of the ITBN component.

|

0.8

!

I

. . . . .

-

A ;

0.6-

w

0.4

0.2

0

I

I

4

8

Composition

'

I

.

12 (ITBN Wt

.

.

.

.

16

~)

Figure 17. The critical stress intensity factor, K c, for the KerimidJTBN blend system as a function of composition at high (curve B) and low (curve A) deformation rate.

152 180

I

,

, I

,I

150 -

E

A

12o -

.,,...

d

9o

60

3O

I

0

I

4

I

8

12

16

Composition (ITBN Wt ~)

Figure 18. The critical strain energy re~ease rate, Gc, for the Kerimid / ITBN blend system as a function of composition at high (curve B) and low (curve A) dejbrmation rate.

I

I

I

3.5

13.

(.~

3.0

LU

2.5

.

0

.

.

.

I

.

.

.

5

.

I

.

10

.

.

I

15

20

Composition (ITBN Wt ~1

Figure 19. The flexural elastic modulus, E, for the Kerimid,TTBN blend system as a junction of composition.

153 The observed enhancement in tougheness is somewhat lower than that obtained by using an engineering thermoplastic like polyetherimide as second component. This illustrates the lower efficiency of the rubbery phase with respect to thermoplastics in toughening densely crosslinked thermosetting materials. In this case the enhancement in the tougheness parameters may be ascribed to the incorporation of flexible rubber chains within the BMI network which render the network more flexible and easier to deform under loading. References

1.

V. Crivello, J. Polym. Sci., Polym. Chem. Ed., 11, 1185 (1973).

2.

I.K. Varma, Sangita, D. S. Varma, J. Polym. Sci., Polym. Chem. Ed., 22, 1419(1984).

3.

J.E. White, M. D. Scaia, Polymer, 25, 850 (1984).

4.

C.E. Browing, "Advanced Thermoset Composites", J. M. Margolis Ed., Van Nostrand Reinhold Co., New York (1986).

5.

T.T. Serafini et al., US Patent 3, 745, 149 (1973).

6.

T.T. Serafini "Status Review of PRM Polyimides", ACS organic Coatings and Plastic Chemistry, 40, 469 (1979).

7.

A.J. Kinloch, S. J. Shaw, Amer. Chem. Soc. Polym. Mater. Sci. Eng., 49, 307 (1983).

.

9.

M. Bergain, A. Combet, P. Grosjean, Brit. Pat. Spec. 1, 190, 718 (1973). H. D. Stenzenberger, US Pat. 4, 303, 779 (1981).

10. H. D. Stenzenberger, W. Romer, M. Herzog, P. Konig, 33rd Int. SAMPE Symp., 33, 1546 (1988). 1 I. E. Plati, J. G. Williams, Polym. Eng. Sci., 1_55,470 (1975). 12. V. Di Liello, E. Martuscelli, P. Musto, G. Ragosta, G. Scarinzi, Die Ang. Makromol. Chem., 223, 93 (1993). 13. J.N. Sultan, F. J. Mc Garry, Polym. Eng. Sci., 13, 27 (1973).

154 14. L. T. Manzione, J. K. Gillham, C. A. Mc Pherson, J. Appl. Polym. Sci., 2__66, 889(1981). 15. A.K. St. Clair and T.L. St. Clair, Int. Adhesion Adhesives, 1, 249 (1981). 16. I.K. Varma, G.M. Fohlen, J.A. Parker, D.S. Varma, in Polyimides, K.L. Mittal, Ed., Plenum, New York, 1984, vol. 1, pp. 683-694. 17. A. J. Kinloch, S. J. Shaw, Int. Adhesion Adhesives, 5_, 123 (1985). 18. Shinji Takeda, Hiroshi Kakinchi, J. Appl. Polym. Sci., 35, 1351 (1988). 19. P. Kavacic, R. W. Hein, J. Am. Chem. Soc. 81, 1190 (1959). 20. M. Abbate, E. Martuscelli, P. Musto, G. Ragosta, "Toughening of a Bismaleimide Resin by a Maleimido-terminated Liquid Rubber", submitted to J. Appl. Polym. Sci. 21. D.O. Hummel, K. U. Heinen, H. Stenzenberger, H. Siesler, J. Appl. Polym. Sci., ~

2015 (1974).

22. C. Di Giulio, M. Gautier, B. Jasse, J. Appl. Polym. Sci., 291771 (1984). 23. S. F. Parker, S. M. Mason, K. P. J. Williams, SpectTochimica Acta, 46A, 121 (1990). 24. Silas, Yales, Thornton, Anal. Chem., 31,529 (1959).

PART II TOUGHENED THERMOPLASTICS

This Page Intentionally Left Blank

157 CHAPTER 4

NUCLEATION PROCESSES IN TOUGHENED PLASTICS A.Galeski l, Z.Bartczak ~, E.Martuscelli 2 ICentre of Molecular and Macromolecular Studies, Polish Academy of Sciences, 90-362 Lodz, Poland

2Istituto di Ricerca e Tecnologia delle Materie Plastiche, CNR, Via Toiano 6, 80072 Arco Felice, Italy

I. Introduction

The properties of a polymer can be extensively modified by a physical mixing with another polymer. The properties of a resulting polymer blend depend on the composition and processing and also on the physical state of each component at the temperature of application of the blend. Over the years most theoretical and experimental investigations of properties of polymer blends have concerned the systems containing amorphous components (e.g.[ 1]), although blends with crystallizable components are receiving increasing attention (e.g.[2]). Below melting points of components the blends with crystallizable polymers constitute heterogeneous systems because the components separate from each other during crystallization. However, the miscibility of the remaining amorphous phases of the components may be still possible. The incompatibility induced by crystallization of components affects the mechanical properties of the blends,

158 causing frequently their deterioration. On the other hand the mechanical properties of blends are determined also by the properties of crystalline phases, including the overall crystallinity, crystalline morphology and the sizes of crystallites and their aggregates such as spherulites. The average spherulite size in blends is a very important factor influencing their mechanical properties e.g. the yield stress and ultimate strength of the material [3]. The size of spherulites is controlled mainly by the process of primary nucleation. The primary nucleation behavior depends on both the material properties and thermal treatment i.e. thermal conditions for processing and crystallization. In the case of blends the primary nucleation depends also strongly on the blend composition. Primary nucleation behavior in blends has been recently widely studied. Below a brief survey of habits and nature of nucleation in polymers is given and the results of investigations of prhnary nucleation of spherulites in polymer blends containing crystallizing components are summarized.

2. Primary nucleation in polymers Crystals in polymers are grown from nuclei rather than formed uniformly over the entire volume of the material. In the primary nucleation phenomenon in polymers three paths for nucleation can be distinguished: (i) homogeneous nucleation which takes place if no preformed nuclei or foreign surfaces are present, (ii) heterogeneous nucleation- the nuclei are formed on foreign surfaces which often reduce the nucleus size needed for stable crystal growth, and thus enhance the nucleation process and (iii) specific only for polymers the self-seeding - the nucleation caused by small polymer crystals which survived melting or dissolution of the polymer sample [4]. The classical concept of crystal nucleation based on the assumption that fluctuations in the undercooled phase can overcome energy bmxier at the surface

159 of the crystal was first developed by Gibbs and later by Kossel and Volmer (see general surveys of nucleation by Zettlemoyer [5] and by Price for macromolecules [6]). The rate of nucleation I* has been derived by Turnbull and Fisher [7] to be

I*=(NkT/h)exp[-(AG*+AG~)/kT]

(1)

where N is related to the number of crystallizable elements, AG~ is the energy of formation of a nucleus of critical size and AG~ is the activation energy for chain transport. The formula derivation is based on the above assumptions using the absolute rate theory. Generally, in polymers as the temperature is lowered from the melting temperature a rapid decrease in AG* and a slow increase in AGn occur causing I* to increase. As the temperature is lowered even further, the decrease in AG* becomes moderate but the increase in AG~ more significant which result in a decrease in I*. Therefore, a maximum in I* exists which is related to the ease with which crystallizable elements can cross the phase boundary. The theories of polymer crystallization are still considerably controversial. The behavior of polymer melts on the molecular level at melting temperature and below it is not yet fully understood and creates problems with interpreting crystallization on higher levels. The basic problem lies in quantification of processes which are only qualitatively understood. For this reason one will fmd in the literature a variety of expressions for a given parameter for polymer nucleation and crystallization. Although the expressions may only differ slightly one has a choice in selecting an expression to use for his own data. The calculations may result in 10-20 % or greater departure from the published data. Since the smallest value of AG* is related to the size and shape of the nucleus in such a way that it has the minimum surface free energy therefore the critical dimensions of the nucleus can be calculated for the anticipated geometry of the

160 nucleus. The expressions for critical sizes of nucleus can be obtained by zeroing the first derivative of AG* (the sum of changes in bulk and surface energies of the nucleus) with respect to the dimensions of the nucleus. Similarly several various expressions can be found for the free energy of fusion Agf. Together with various geometries of the nucleus it gives rise to a range of equations which relate important parameters. Hoffman [4,8-10] presented extensive diverse work in this area. For this reason the expressions derived or used by Hoffinan will be quoted in this review. For example for the case of homogeneous primary nucleation and for rectangular shape of the nucleus the free enthalpy of the formation of nucleus of critical size can be described by the following expression (e.g. see Ref.[4])

AG*=(32(yrr~)/(Agr)2=[32~zco(Tm)2]/[(Ahf)2ta(AT)2]

(2)

where ~ and r are the side and end surface free energies of the crystal,

Agf is the

free enthalpy of fusion of the crystal of the chosen geometry, T ~

is the equili-

brium melting temperature, AT= TO-T, and f=2T/(T ~ +T). The term AGn is usually approximated by the WLF equation for the viscous flow:

AGn/kT=U*/[R(T-Too)]

(3)

Based upon these types of calculations, one will find that the typical homogeneous nucleus dimensions are about 103 to 105 A 3 while a typical polymer chain voltune is about 105 to 107 ,~3. Thus, only a small portion of the polymer chain is involved in forming a nucleus. One of the two types of nuclei - fringed micelle - is thought to be a bundle of polymer chains with long sections remaining uncrystallized. There are restrictions to the fringed micelle formation. As it was shown by Flory [11] the strain generated at the crystalline-amorphous interface by polymer molecules which cross the interphase boundary must limit the nucleus and crystal

161 dimensions. A single molecule, therefore, must fold in order to reach the proper dimensions for the formation of a nucleus and fitrther growth. Thus, the crystal continues to grow in chain folded fashion with constant lamellar thickness. Chain folded nuclei are more probable than fringed micelle nuclei in all cases where the segmental mobility of macromolecular chain is high. Another path of primary nucleation is the heterogeneous nucleation studied extensively by Binsbergen [12-17]. The experimental part of his works concern mostly the heterogeneous nucleation in isotactic polypropylene. Based on the vast number of his own and other authors experimental observations Binsbergen derived a theory of heterogeneous nucleation of crystallization in polymers [17]. The formation of a nucleus on a foreign surface involves a creation of a new interface, similarly as in the case of homogeneous nucleation. However, the preexisting foreign surface greatly reduces the free enthalpy of the formation of a critical nucleus, AG*. This lowers the critical size of the nucleus and results in the formation of heterogeneous nuclei at lower undercooling. Again, assuming rectangular shape of a nucleus lying fiat on a foreign surface one can obtain the expression for the free enthalpy of the formation of a nucleus of the critical size:

AG*=( 16Accyce)/(Agf)2=[ 16AacrCre(T~ )2]/(AhffAT)2

(4)

where Aa is the specific interracial free energy difference for the interface: nucleus-foreign surface. Similarly as for the homogeneous nucleation AG* in the equation (4) is proportional to 1/(AT)2. However, for very active foreign surfaces characterized by low value of A~ the critical thickness of a nucleus approaches the molecular thickness, and the formula (4) transforms to a form

AG*=(4bo~gr Tm)/(AhffAT - Ac Tm/bo)

(5)

162 where bo is the molecular thickness. As At~ goes to zero the formula (5) approaches the enthalpy barrier characteristic for secondary nucleation mechanism which is proportional to 1/AT. The kinetic nucleation theory with chain folding provides now the best general tool for understanding the primary nucleation and the growth of polymer crystals at isothermal conditions from unstrained melt [4]. The reptation concept proposed originally by de Gennes [18] was also adapted for the description of chain motion and transport in the melt [8,19,20]. The reptation theory leads to more accurate expressions for the preexponential factor in I~ and for the activation energy for chain transport, AGn, in eq.(1). It also predicts the dependence of the crystallization rate on molecular weight in different regimes of crystallization [9,10]. For example the appropriate expression for the secondary nucleation process predicted by the kinetic nucleation theory with reptation is as follows [8]:

I=(Nol3gpi)/(aon0exp[-4booa~ T~

(6)

where No is the number of reacting species at the growth front, 13g=(~n)(kT/h)exp[Qo - RT], n is the number of macromolecule segments in the melt, ns is the number of stems of width ao, h is the Planck constant, and QD is the activation energy for reptation. K is a constant usually of the order of unity as determined from the experiments. The most spectacular prediction of the reptation concept concerns the mean time of reeling out from the melt an entire macromolecule with one end attached to the nucleus onto the growing from and pulled by crystallization forces [9]:

t= 1.9xlO9n2 [s]

(7)

163 (n is the number of chain units, other required parameters taken for polyethylene) which is very short (order of 10Zs). For comparison, the time for establishing intermolecular entanglements in a polymer melt is approximately four orders of magnitude longer [21] (intermolecular entanglements were removed by dissolution in a solvent followed by quick evaporation of a solvent and careful drying ). The time for restoring the intermolecular entanglements is of order of tens of minutes for polyethylene melt. These two results point out that macromolecules in the melt are relatively immobile in contrast to a crystallizing macromolecule pulled at one end by attractive forces of crystallization. It is also clear in the view of those results that the nucleation of new crystalline layers on an existing crystal is the controlling factor of the crystal growth. Direct evidence for further nucleation step beyond primary nucleation was brought by Wunderlich and Cormier [22] from observation of crystallization of linear polyethylene melt seeded with extended chain crystals. The observable crystal growth is a result of two processes, the first being the nucleation of initiating stems on the surface of the crystal, and the second being the coverage of the surface by new stems beginning at the initial stem. It must be emphasized here that the crystal grows macroscopically in the direction normal to its surface while on the molecular level the elementary growth mechanism is the growth along the crystal surface.

3. General remarks on primary nucleation.

From the expressions for AG* for homogeneous (eq.(2)) and for heterogeneous (eqs.(4) and (5)) nucleations the constant nucleation rate in isothermal conditions is expected. However, in polymer samples there are usually heterogeneous seeds with a broad specmma of Act values resulting in various nucleation rates I*. Also a limited number of those seeds present in samples leads to differentiated

164 exhaustion of particular fraction of nuclei. Moreover, the self-seeding gives rise to an almost instantaneous nucleation. All those attributes of nucleation events cause that the real nucleation process in polymers is a complicated function of time, not just a temperature: I'=I*(T(t),t). The theories of homogeneous, heterogeneous and self-seeded nucleations describe the mechanisms and show the tendency but hardly predict the real habit of nucleation in a given polymers. Hence, the experimental methods of determination of nucleation are of particular importance. The knowledge of nucleation data is often essential for controlling physical properties of polymer; mechanical properties depend to a great extend on the spherulite average size, size distribution and the size of so-called "weak spots" defects of spherulitic structure including cavities and frozen stresses which resulted from volume contraction during crystallization [23,24], all determined by the primary nucleation process. For some applications it is sufficient to determine only the total number of nuclei activated during the crystallization. The simplest way of obtaining this value is from the average spherulite size for samples which are filled with spherulites. The average spherulite size can be obtained from the first moment of a size distribution or of a distribution of chord intercepts with spherulite boundaries [25]. Other average spherulite sizes can be obtained from higher moments of spherulite size distribution. The higher moments of spherulite size dislaibution can be determined on the basis of direct characterization of spherulite patterns (under polarized light microscope, under scanning or transmission electron microscope of thin films or sections of bulk samples - second or third moment of the spherulite size distribution), on the basis of the small angle light scattering (fourth or fifth moment) (e.g.[26,27], also [28]) or of the light depolarization technique (second or third moment) [29,30]. However, if the time dependence of activation of nuclei is required other methods must be used. The data on time distribution of primary nucleation are usually

165 obtained by direct microscopic observation of a crystallizing sample. The odds of this method are the necessity of using thin samples and the condition of crystallization allowing for counting the spherulite centers. Those limitations can be overcome by applying a method of reconstructing the sequence of nucleation events from shapes of spherulite boundaries in already crystallized films [31] and in thin sections for bulk samples [32]. The time lag between the nucleation of two neighboring spherulites can be found from the curvature of their common boundary and this procedure repeated for a chain of neighboring spherulites delivers the data on the time distribution of the activation of nuclei. The time distribution of activation of nuclei should be expressed in number of activated nuclei per volume unit of untransformed fraction of the sample. Calibration of the time axis in that method is made by measurements of the spherulite growth rate. Usually, a given brand of polymer is characterized by an average number of primary nuclei per volume unit at certain crystallization conditions. The average spherulite size in bulk is determined by the number of nuclei per volume unit. For thin films, however, the spherulites as seen in plane are larger. For the thickness of a sample below the average spherulite size in bulk, the thinner the sample the larger the spherulites. The apparent increase in spherulite sizes in thin films results from the constant average number of nuclei per volume unit. The factor complicating this simple relation is the nucleating ability of sample surfaces. The Avrami type of analysis is often erroneously applied for obtaining the nucleation data from differential scanning calorimetry isothermal crystallization experiments and from dilatometry. The reason for this is that the conversion of melt to spherulites is assumed to follow pure sporadic or pure instantaneous modes, ~e .only modes described correctly by. the Awami equation"

ct(t)= 1-exp(-Kt" )

(8)

166 where c~(t) is the degree of the conversion of melt to spherulites. However, the general form of the equation for conversion kinetics of melt to spherulites for isothermal experiments, which was first developed by Avrami is as follows:

a(t)- 1-ex(-pG2

a(t) - 1 - ex

fo I ( t ) ( t - t ) 2 dt I

- -~ pG 3

(t)(t- 0 3 dt

for films

(9a)

for bulk samples

(9b)

where G is the spherulite growth rate constant at a given temperature of crystallization and I(t) is the rate of nucleation. It is evident that eqs.(9a) and (9b) that for I*=IoS(t) (5(0 being the delta Dirac function; this represents the instantaneous mode of nucleation) one obtains in the 'Avrami' exponent t2 and t3 *

c

for two and three dimensional cases, respectively, while for I =Iot, (c being an integer number, sporadic nucleation for c=0) the 'Avrami' exponent is proportional to tr and tr for the two and three dimensional cases, respectively. Therefore the plot of ln[-ln(c~(t))] vs. ln(t) as required for the Avrami type of analysis can lead to a straight line only in two limiting cases: instantaneous and sporadic nucleation modes, both very seldom in a plain form in crystallization of polymers. The proper way of data analysis is by solving the integral equation which follows from eqs.(9a) and (9b)

~2 I(t)(t- t) u dt = - 1 / ( p G u)ln[1- a(t)]

for films

(10a)

~~ I ( t ) ( t - t ) 3 d t = - 3 / ( 4 p G 3)ln[1 - a(t)]

for bulk samples

(10b)

167 The simplest way of solving it is by Laplace transformation, or equivalent, by third (for films) or forth (for bulk) order differentiation against time, t: I(t) = -1/(2pG 2)d3 [ln(1- a(t))]/dt 3

for fihaas

(1 la)

I(t) = -1/(8pG 3)d4 [ln(1 - a(t))]/dt

for bulk samples

(11b)

4

In the literature there are many examples of the treatment of the problem of nonisothermal solidification ( for a broad review see the paper by Wasiak [33]). Most of them are lacking in a firm theoretical background. However, the probabilistic approach to the description of spherulite patterns [34,35] provides now a convenient tool for the description of the conversion of melt to spherulites. In the case of nonisothermal crystallization the conversion of melt to spherulites is described by [34,35 see also 36]:

a(t) = 1-exp{-pf~ I(t)If~ G(s)ds] 2 dt}

for films

(12a)

a(t) = 1-ex~-(4p/3)~~ I(t)I~~ G(s)dsl 3 dt}

for bulk samples

(12b)

Since the growth rate is unambiguously determined in all three regimes of

crystallization by secondary nucleation process and completion rate of the nucleated layer it could be precisely meast~ed in isothermal experiments in thin films as a function of temperature and could be then easily transformed to the function of time provided that the temperature change is monitored during

168 nonisothermal solidification of a polymer. The solution of the eqs.(12a) and (12b) is also by Laplace transformation, or equivalent, by differentiation [36]:

d { 1/G(t)-~d I 1/G(t)-d-d~-{ln[1-a(t)]}l I for films I(t)= -1/(2P)-d-~

{

dE d{

(13a)

dI,n ,a 0 jlltforb mp,es ,3b

I(t)=-l/(8p) d 1/G(t)-d-~ 1/G(t)-d-~ 1/G(t)~

In all above equations the conversion degree of melt to spherulites must not be mistaken for the crystallinity degree. The difficulty in all nonisothermal experiments is that the fractions of the material crystallized at different temperatures differ in the degree of crystaUinity. It is usually considered that nuclei are spread randomly over the sample, except for nuclei formed on outer surfaces in three dimensional samples, on surfaces allowing for transcrystallinity and on surfaces of the second dispersed component e.g. short fibers. However, that is not necessarily true. If the nucleation events occur not only at the very beginning of crystallization but also later during crystallization process then the volume occupied by already advanced spherulites are excluded from further nucleation. The close vicinity of an arbitrarily chosen nucleus is then poorer in other nuclei than more distant regions. Although the nucleation itself is a spatially random process it is limited to the uncrystallized portion of the sample. Such an exclusion always produces a kind of a distance correlation, if the nucleation process is prolonged in time. The spatial correlation of spherulite centers was first observed by Misra, Prud'homme and Stein [37] while respective mathematical formulas for the description of the distance correlation of nuclei for model modes of primary nucleation were derived in Ref.[35] and [36] for isothermal and nonisothermal cases.

169 4. Primary nucleation in polypropylene.

Isotactic polypropylene is now very frequently used as a base for many blends. The world production of polypropylene exceeded in 1986 8 million ton and about 60% of this production went into blends with other polymers. These numbers show an increasing tendency (e.g.[38]). Therefore, the nucleation habits of isotactic polypropylene, being very important example of crystallizing component of polymer blends, are revealed below. Polypropylene has been shown to exhibit several crystalline modifications [39-42] in addition to the most common ot monoclinic structure reported first by Natta et al.[40]. The [3-phase crystallizes from primary nuclei in the hexagonal fashion, although, the [3-phase can also be initiated along a growing front of the a-phase in the temperature gradient. The T triclinic phase was first discovered in low molecular fractions of isotactic polypropylene solidified by slow cooling [41], in commercial polypropylene crystallized under high pressure [42] and in samples of propylene-ethylene copolymer at a low ethylene content which exhibited a peculiarity of a complete crystallization in the T-form [43,44]. In the literature there are no reports on primary nucleation of the T-form spherulites, however, the 7-form is reported to exhibit spherulitic lamellar structure and undergoes the y-or transition on annealing at 147~ and 1 atm [45].

4.1 Nucleation of a form of isotactic polypropylene.

First consistent microscopic data on primary nucleation of isotactic polypropylene were obtained by von Falkai and Smart [46] and also by von Falkai [47]. Their data for samples melt annealed at 180~

prior to crystallization and then

crystallized isothermally in the temperature range from 122 to 145 ~ presented in Table I.

are

170 Table h Basic crystallization parameters o f isotactic polypropylene. ~

Crystallization

Nucleation

Growth rate

temperature [~

density [ 106 cm3]

122.0

2.36

18.0

125.0

1.56

12.0

127.5

1.02

7.0

130.0

0.85

4.3

132.5

0.73

2.6

135.0

0.65

1.6

138.0

0.58

0.86

140.0

0.53

0.59

145.0

0.47

0.27

[taroJmin]

~Data from von Falkai and Smart [46] for isotactic polypropylene of Mw=51 200, heptane extracted, melt annealed before crystallization at 180~ for 15 min

The nucleation in these experiments in polypropylene was found to be heterogeneous and instantaneous with calculated Avrami exponent closely matching 3. (There is somewhat unclear point about the three-dimensionality of their samples used for microscopic observation of crystallization). Since then many authors published nucleation and crystallization data for isotactic polypropylene (e.g.[29,30,48-52], see also [53]). The change from instantaneous to sporadic character of primary nucleation is reported in the literature when the samples were heated up to 200~ and above

171 [48,51]. Under such condition the polarized light microscopy examination revealed an initially constant nucleation rate, however, decreasing for longer crystallization time. The Avrami type of fit to integral exponent was not satisfactory in this case. The course of primary nucleation in isotactic polypropylene down to 70~ was first demonstrated by Burns and Tumbull [54] and by Koutsky, Walton and Baer [55] employing the droplet technique developed originally by Vonnegut for tin and water droplets [56]. One can recognize four distinct regions in the nucleation of isotactic polypropylene melt: (i) immediately below the DSC determined melting point (165-167~

there is a

gap where the crystal nucleation and growth hardly takes place. Neither the present heterogeneities nor introduced nucleating agent can accelerate the nucleation. (ii)Most of published nucleation data concern the region of temperature below ~50~

but above 115~

where regular spherulite are nucleated (see

e.g.[57,58]) (although Binsbergen and deLange [59] observed negatively birefringent sheaves of crystalline lamellae crystallized isothermally at as high as 160~ which is well above Regime II - Regime I transition temperature estimated for isotactic polypropylene at 155~

[58]). This region is the

extended region of activity of heterogeneous nuclei. The number of heterogeneous nuclei is limited during the crystallization. (iii) Some of the heterogeneous nuclei become active at even lower temperature which follows from their smaller size or lower perfection. These nuclei are also limited in nmnber. (iv) Finally at approx.80-85~

and below there is the region of homogeneous

nucleation. The number of nuclei in this region increases rapidly with the decrease of the temperature.

172 Except for very thin specimens, it is difficult to reach beyond the upper range of activity of heterogeneous nuclei due to the low thermal diffusivity of polymers, intense nucleation and fast spherulite growth at lower temperature. Also the latent heat of fusion liberated by the rapid crystallization during quenching tends to maintain the temperature during crystallization in the upper range of activity of heterogeneous nuclei. This and the instantaneous character of most of heterogeneous nuclei cause that homogeneous nucleation range is rarely reached and many of polymeric objects in technological applications crystallize only from heterogeneous nuclei. Figure 1 illustrates the nucleation activity in isotactic polypropylene (RAPRA, iPP1, Mw=3.07*105, Mn=l.56*104.. density= 0.906 g/cm 3, melt flow index=3.9 g/10 min). The data were taken from Refs.[60-62] and differentiated to represent the contribution of new nuclei activated by the temperature decrease by 1~ (data for the crystallization temperature of 90, 100 and 110~ were obtained by the authors for the purpose of this review employing the method of crystallization described in Ref.[62]). The smnples in a fonn of thin films (20-30 ~ n ) were crystallized isothermally on a microscopic hot stage for the temperature above 115~ while for the temperature below 115~ the samples were obtained by isothermal crystallization in a specially designed crystallization cell enabling to reach isothermal conditions within the sample volume in less than 0.5s. It is seen that the number of nuclei increases initially as the temperature of crystallization decreases. At the temperature of crystallization of 132~ a change of slope of the AI/AT vs. T curve is seen which is apparently associated with the regime II- regime III transition in crystal growth kinetics (reported to be at 135 - 137~ for other brands of isotactic polypropylene as determined in Ref.[10] on the basis of the data taken from Refs. [12], [46] and [63], see also Ref. [58]). The regime IIregime III transition is in fact the change in the intensity and the habit of secondary nucleation which may be considered as heterogeneous nucleation on the polymer

173 crystal surface; similar though not identical transition (slightly different temperature of the transition) should be expected for nucleation on surfaces of other heterogeneous nuclei. Further decrease of the crystallization temperature below 115~ results in saturation of the AI/AT value. Apparently all heterogeneous nuclei present in the sample are able to show up within the time of crystallization below 115~ At the temperature below 85~ a new intense process of homogeneous nucleation takes place. A rapid increase in number of formed nuclei with the decrease of crystallization temperature is observed (see Fig. 1). Early droplet experiments during isothermal crystallization [54,55] also showed that the nucleation in droplets of isotactic polypropylene is thermally activated and the droplets crystallize sporadically in time. Investigations of nucleation performed during continuous cooling could, however, resolve only a singular large peak of nucleation at one particular undercooling. Annealing of the melt has a great influence on primary nucleation in isotactic polypropylene. However, the knowledge of the behavior of primary nucleation during melt annealing in polypropylene was acquired gradually as the understanding of nucleation processes in polymers became better. First extensive study of the effects of thermal history on crystallization of isotactic polypropylene was conducted by Pae and Sauer [64] and Sauer and Pae [65]. Further studies included direct microscopic observations of the formation of spherulites in polypropylene melt subjected to various thermal treatment. Annealing of polypropylene melt prior to crystallization decreases the active fraction of primary nuclei. The crucial factor is the temperature of melt annealing below or above the equilibrium melting temperature, T ~ . At 190~ a vast number of those thermally sensitive nuclei remains untouched while even short exposure to temperature around 220~ decreases the number of active nuclei by orders of magnitude.

174 107

"

L

~

.

.

.

.

.

.

~

'

[

'

10 6 05

10 4 I.....I

[.~
60~

while it is immiscible with

PMMA. As a consequence, the growing of a PMMA phase can, in principle, give rise to a complex morphology with the separation of a minor phase

450 constituted by EVA droplets trapping some PMMA particles.

The final

microstructure that then might develop, schematically represented in Figure 6, can be described as a multicore shell structure, very similar to that industrially realized. We have demonstrated that the addition of as much as 7% of rubber is sufficient to cause a high improvement of the impact properties without a significant loss of tensile modulus, while the optical transparency, typical of methacrylic matrices, is maintained in most of the blends [19-21]

Fig. 6. Schematic representation of the multicore~shell PMMMEVA synthetic blends.

structure in

2. Synthesis of Blends

The blends were obtained by the following simple method: the EVA copolymers were dissolved in MMA monomer in a ratio 7/100 by weight. After adding a 0.2% by weight of benzoyl peroxide as radical source, a first stage of polymerization was effected under efficient stirring. When the viscosity reached a critical level (normally after about 100 min.), a final curing step was effected in

451 a mold kept at 80~ for 12 h. It is worth noting that the MMA monomer (Fluka product) has not been purified from the inhibitor.

Table H

Codes, composition, and rl.,Y of the used EVA copolymers

Code

VA % by weight

(dL/g) LMWl8

18

0.54

LMW28

28

0.54

MMW40

40

0.70

MMW20

20

0.83

HMW20

20

1.04

HMW9

9

1.06

HMW10

10

1.30

HMW36

36

1.38

~)c = 0.25 g/(100mL toluene) at 30~

As can be observed in Table II, the employed EVA copolymers, kindly supplied by Dupont, show a VA content ranging from 9 to 40% by weight with rlmh values ranging from 0.5 to 1.4 dL/g (viscosimetric analysis was effected in toluene by a Ubbehlode viscosimeter). molecular weight

The employed codes

refer to low

(LMW), medium molecular weight ( M W ) ,

and high

molecular weight (HMW) EVA, followed by the VA content (% by weight). The same codes are hereafter used to identify the prepared blends. The plain PMMA used as reference was prepared by a conventional radical process initiated by organic peroxide (0.2 wt %).

452

3. Morphological Analysis

Electron microscopy spectroscopy was used to characterize the final morphology of the blends and to investigate on the phase inversion process and on the fracture mechanisms. A Scanning Electron Microscope (SEM), Philips 501 model, was used. The analysis has been performed on fracturated surfaces, coated with a thin layer of gold/palladium alloy. Some blends were subjected to a smoothing and polishing procedure and observed after exposure to n-heptane vapors (20 min) to remove the EVA phase.

3.1 Phase inversion

To study the morphological development of phases in our blends, we observed smoothed and polished surfaces after n-heptane extraction. The final morphology of a blend having medium-viscosity EVA is reported in Fig. 7. It is evident that the morphology revealed by the extraction of EVA is very complex. Particularly, the PMMA constitutes the matrix while large regions of EVA with subincluded PMMA particles are deafly evident. High viscosity EVA blends show the same final morphology depicted in Fig. 7. The revealed morphology is very similar to that reported for HIPS [4,5]. For HIPS, it also reported that the absence of stirring leads to a "not inverted" morpholgy, in which a matrix of rubber surrounds large spherical polystyrene domains even if the rubber represents only 7% of the total polymeric material. This situation was obtained also by us and is shown in Fig. 8 where the same blend of Fig. 7 was polymerized under stirring for only 90 mm, i.e., before the attainment of the proper viscosity for "phase inversion". Upon etching with n-heptane, the tiny network of rubber that surrounded PMMA particles is removed. It is evident that PMMA, altough present in a proportion larger than

453 90%, nevertheless constitutes the dispersed phase embedded in a continuous network of EVA copolymer.

Fig. 8. SEM micrograph of a smoothed surface of MMW20 blend (90 min of stirring) aRer etching with n-heptane (640x)

454 The above results deafly show that a so-called phase inversion process must occur during the polymerization. The final morphology then strongly depends upon the stirring conditions. For low-viscosity EVA, even after 110 mm of stirring, the morphology of the cured blend, reported in Fig. 9 after smoothing and etching with n-heptane, shows a "not reverted" situation, with large domains of PMMA surrounded by a tiny shell of EVA. This demonstrates that the MW of the rubber also plays an important role in the development of the final morphology.

Fig. 9. SEM micrograph of a smoothed surface of LMWl8 blend after etching with n-heptane (640x)

3.2 Fractographic Analysis The fractographic analysis of PMMA and PMMA blends surfaces after impact testing at room temperature is reported in Figures 10-13. For all the figures, the notch front is on the left-hand side.

455

Fig. 10. SEM micrograph of fracture surface at room temperature of plain PMMA (320x) The fracture surface of PMMA (Fig.10) shows a series of brittle fracture bands or striations, oriented perpendicular to the crack propagation.

The

formation of these bands is because, above a certain crack speed, the craze preceding the crack front undergoes branching. [22]. At stn~ciently high stress levels, these crazes undergo fracture, causing surface roughening (bands), a deceleration of the crack, and a drop in the stress amplitude around the crack tip of very large domains in a matrix of a tiny layer of EVA copolymer. A similar structure, as already mentioned, is consistent with the absence of phase inversion during the radical polymerization of MMA.

Apart from this observation, the

fracture surface does not reveal any fractographic feature that is below that necessary to initiate branching crazes. The fracture then reverts back, the crack spee~ rises again, and branching reoccurs. The repetition of this process gives rise to the banded appearence of the fracture surface.

Figure 11 shows the

fracture morphology of a blend containing low viscosity EVA copolymer. This blend is characterized by a major-phase PMMA dispersed in the shape distinctive

456 of plastic deformation mechanisms (crazes and/or shear bands) from which crack development can take place.

Fig. 11. SEM micrograph of fracture surface at room temperature of LMWl 8 blend (320x) Completely different are the fracture surfaces of blends containing medium and high-viscosity EVA copolymers, as can be seen from the micrographs of Figures 12-13, respectively. The overall morphology of these blends seems to be very similar. In fact, both photos reveal domains (less than 1 micron in size) finely dispersed and well embedded in the matrix.

Moreover, signs of an

extensive plastic deformation in the matrix are also evident.

In view of the

intrinsic morphologies revealed by the heptane extraction (see previous paragraph), we may attribute the small particles of Figs. 12 and 13 to the glassy PMMA particles subinduded in the rubbery domains, these last domains being highly deformed during the fracture. This indicates that a large amount of energy is dissipated, probably in the form of crazes, during the impact process.

The

above considerations account for the very high impact toughness observed in such materials.

457

Fig. 13. SEM micrograph of fracture surface at room temperature of HMW20 blend (320x)

458 The influence of the VA content of EVA copolymers on the impact properties of blends can be explained as follows: for low molecular weight EVA, where, as we have previously reported, the phase reversion process does not occur during the stage of polymerization under stirring, an increase in the VA content is to improve the polarity of EVA copolymers and to create a stronger interracial adhesion between EVA and PMMA with a consequent improvement in teh impact properties. For MMW blends, such effects are less evident. No large differences are observed in the impact properties as function of VA content of EVA rubbers, at least at room temperature.

4. Mechanical properties 4.1. Impact Behavior

The impact properties were analyzed according to the Linear Elastic Fracture Mechanics (LEFM) approach [23].

The procedure used for the

calculation of the critical strain energy release rate (Gc) and the critical stress intensity factor (Kc) is reported elsewhere [24]. Charpy-type specimens (6.0 mm wide and 60 mm long) were cut by a mill and notched with a fresh razor blade. Then, they were fracturated at different temperatures and at an impact speed of 1 m/s by using an instrumented Charpy pendulum. To obtain sheets of 3.00 mm thickness, the blends and the pure PMMA were compression-moulded in a heated press at 200~ and at a pressure of 240 atm. For simplicity, the impact values can be collected into three groups as function of the molecular characteristic of employed EVA, i.e.: (a) Blends containing low-viscosity EVA (rl~ = 0.54 dL/g). (b) Blends containing medium-viscosity EVA (rl~ = 0.7-0.83 dL/g). (c) Blends containing high-viscosity EVA (rl~ = 1.04-1.38 dL/g).

459 In Figures 14-16, the values of Gc (energy release factor) vs. temperature are reported for, respectively, low-, medium-, and high-viscosity EVA.

Gc (KJ/m=)

2.0

O LMW 18 VA 18% ~inh 0.54 9LMW 28 " 28% .... 0.54

1.5

J

1.0

-~-'~"~

~

PMMA

0.5[

-80

-60

-40

-20

0

T(~

Fig. 14. Energy-release factor Gc of plato PMMA and of PMMA/lowviscosity EVA blends as function of the testing temperature

Gc( KJ/rr~ 3.0

*

/

MMW40

9 MMW 20

VA 40% ~ i n h 0.70

......

2.0.

f

1.0~ v

9

-80

-at

A

.

.

v

&

-60

,

v

I

-40

.

.

v

I

-20

.

v

I

0

v

PMMA

I

I

20

40

T(*C)

Fig. 15. Energy-release factor Gc of plato PMMA and of PMMA/medium viscosity EVA blends as function of the testing temperature.

460

G ( K J / m =) 3.0

9 H M W 20 VA 20% ~ i n h zx H M W 36 " 36% . . . . o H M W 10 ,, 10% . . . . 9 HMW 9 9% . . . .

1.04 1.38 1.30 1.06

2.0

PMMAIPVAC 1.0

-80

PMMA

-60

-40

-2

0

20 T(~.)

Fig. 16. Energy-release factor Gc of plato PMMA, of PMMA/highviscosity EVA, and of PMMA/PVAc blends as function of the testing temperature. As can be observed from the impact data, the blends obtained with lowviscosity EVA are characterized by Gc values similar to, or worse than, that of pure PMMA, in all the ranges of investigated temperatures. Going to medium-viscosity EVA (see Fig. 15), the behaviour of blends is definitely more satisfactory.

In fact, the values of Gc for both kinds of

copolymers are well above the values of PMMA for all the investigated temperatures. The major improvement is observed at temperatures higher than 10~

where PMMA still behaves as a brittle material, while the blends undergo

brittle-to-ductile transition. The differences in the Gc values of the two blends are relevant at temperatures below 0~

while their curves tend to overlap at

room temperature. A similar trend is observed for the blends reported in Fig. 16 having EVA copolymers with the highest molecular weights. Another important variable that we have studied is the vinyl acetate (VA) content. In Table 111, the Gc values at 20~ are reported. For blends with low MW EVA, the Gc diminishes, increasing

461 the VA content, but at high MW EVA (i.e., rl~ > 1 dL/g), the Gc goes through a maximum.

Table Ill

Codes of the blends, % of VA content of the rubber phase in each blend, and critical strata of energy release rate (Gc).

Code

VA content (%wt)

Gc (20~

LMWl8

18

1.3

LMW28

28

0.9

MMW40

40

3.4

MMW20

20

3.5

HMW20

20

3.6

HMW9

9

2.3

HMWl0

10

2.5

HMW36

36

3.2

PMMA

-

1.0

PMMA/PVAc

100

1.1

4.2. Tensile properties

As a high elastic modulus is an important characteristic of PMMA, tensile testing was performed in order to investigate about the amount of its forecastable diminution in these new rubber-modified PMMA based blends. We tested the

462 plain PMMA used as reference and the MMW20 blend, as it showed complete phase-inverted

morphology,

excellent impact properties

and

outstanding

trasparency. Tensile testing was performed using an Instron machine. Specimens of 1.6 9 4.5 930 mm were obtained from sheets compression-moulded at 200~ heated press at 240 atm. The samples were tested at 20~

in an

and at a constant

cross-head speed of 10 mm. mm -~ . From the stress-strata plots, the values of Young elastic modulus (E) and of ultimate tensile strength ( ~ ) were obtained (Table IV). It can be observed that the diminution of E is neglectable, as consequence of the low weight content of the rubber.

On the contrary, the ~

value shows an abrupt diminution,

comparable to those observed in classical PMMA-rubber toughened materials. This diminution is a consequence of the high apparent volume ratio between dispersed phase and matrix, due to the core-shell structure which develops in these systems.

Table IV Young's elastic modulus (E) and ultimate tensile strenght ( ~ ) of plain PMMA and MMW20 blend Code

E (Kg. cm -:)

cr~ (Kg. cm -~)

PMMA

1.5.104

440

MMW20

1.3 9104

250

5. Thermal Analysis The thermal properties of EVA rubber and of a PMMA-EVA blend (namely, the MMW20 blend) have been analyzed in the DSC experiments

463 reported in Figs. 17-18, respectively. A Mettler System TA-3000, equipped with a control and programming unit (microprocessor TC 10A), was used. The system was provided with a calorimetric cell DSC-30, which allowed temperature scans from -170 to 600~

The experiments were performed under

a nitrogen flow. Fig. 17a shows the DSC melting trace for EVA MMW20 and Fig. 17b shows the crystallization trace. The EVA rubber has two clear melting peaks (at about 50~ and 90~

and two clear crystallization peaks (at about 65~

and

40~

1"

1"

E

o_ E ,-

.u

,-

-,--1 ~ - ~

-100

T"--r"-'-T

-50

0

~'-"-~'-"--r --'-"-"

50

Temperature

-1-~" " - " - "- I . . . .

100

150

1 ~

200

9~ " - T ' -

250

2so

2oo'-';Co

,~o

Temperature

(~

(a)

~'o

o

.so

-,oo

(~

(b)

Fig. 17. DSC traces of EVA" (a) meltmg~ and (b) crystallization.

Figs. 18a and 18b show DSC traces for the melting and crystallization, respectively, of the MMW20 blend. crystallization peaks between 50~

The blend shows weaker melting and and 90~

Fig.18 also shows a transition at about 110~ transition of the PMMA phase.

and between 70~

and 25~

this corresponds to the glass

464

1"

1'

u E o

i.u

Temperature (~

(a)

Temperature (0C)

(b)

Fig. 18. DSC traces of MMW20 blend during" (a) melting, and (b) crystallization.

6.Graft copolymers formation The high impact properties of HIPS are at least partially interpreted as a result of the formation of graft copolymer species between polybutadiene and the growing chains of polystyrene. In our case, we used radical polymerization to grow PMMA, and EVA copolymers are constituted by long polyethylene sequences that are known to be reactive toward radical copolymerization. To get indirect proof that grafted EVA-g-PMMA species are formed in our system, we prepared and characterized a blend using poly(vinyl acetate) as a second phase. The Gc values at different temperatures reported in Figure 16 and the value of Gc at room temperature reported in Table 111 show that this blend has very poor impact properties, comparable with those of plain PMMA.

465 Also, the fracture surface (see Fig 19) resembles that of unmodified PMMA.

The surface is covered with brittle fracture bands originated from a

cyclic process of formation and breakdown of crazes.

Fig. 19. SEM micrograph of fracture surface at room temperature of PMMA/PVAc blend (160x). We believe that this result indicates that, in the absence of graft copolymerization reaction between PMMA and the dispersed phase, an effective toughening process cannot occur, although the dispersion of the minor phase is very intimate.

This could also explain why the impact properties of HMW

blends seem to go through a maximum as a function of VA content (HMW20>HMW36). In fact, we can suppose that a balance occurs between the increase in interracial compatibility between EVA and PMMA by increasing the VA content (better impact properties) and the decreased ability to form graft

466 copolymer species, due to the diminution of the length of polyethylenic sequences on EVA chains at higher VA content.

7. Optical properties The

PMMA/EVA

blends

which

show

a

complete

phase-inverted

morphology keep a transparency at room temperature very close to that of pure PMMA. This is due to the effect of equalization of the refractive indices of the two components, related to the particular multicore structure. In fact, a blend prepared simply by melt-mixing preformed EVA and PMMA polymers in the same weight percentage (i.e., 7/100 w/w) in a Brabender-like apparatus is completely opaque at any temperature. Moreover, these blends exhibit a very peculiar and interesting optical behavior [19-21].

In fact, altough almost completely transparent at room

temperature, their transparency gradually decreases increasing the temperature until, at T > 70~

they become completely opaque. This temperature-dependent

opacification, which is reversible, has been ascribed to the melting of ethylenic sequencies of the EVA component, which gives rise to a light-scattering phenomenon; taking into account the microstructure of the blend, the description of this kind of scattering can be carried out in the Rayleigh-Gans approximation [251. In conclusion, the optical behavior of these blends appears attractive in view of their possible applications as a temperature-controlled optical device or as an active element for thermally induced optical bistability.

References 1) C.J. Hooley, P.R. Moore, M. Whale, M.J. Williams, Plast. Rubber Process. Appl. k 345 (1985)

467 2) C.B. Bucknall, J.K. Partridge, M.V. Ward, J. Mater. Sci.19, 2064 (1984) 3) L.H. Sperling, "Interpenetrating Polymer Networks and Related Materials", Plenum Press, New York (1981) 4) Yu.S. Lipatov and L.M. Sergeva, "Interpenetrating Polymeric Networks", Naukova Dumka, Kiev (1979). 5) D. Klempner, Angew. Chem. 90, 104 (1978) 6) D.L. Siegfried, D.A. Thomas and L.H. Sperling, J. Appl. Polym. Sci., 26, 177(1981). 7) J.K. Yeo, L.H. Sperling and D.A. Thomas, Polym. Eng. Sci., 21,696, (1981) 8) D. Klempner and D.C. Frisch, eds., "Polymer Alloys 11", Plenum Press, New York (1980). 9) G.E. Molan, J. Polym. Sci., A-3, 4235 (1965) 10)C.G. Bucknall, "Toughening Plastics", Applied Science, London (1977) l l)K. Kato, Polym. Eng. Sci., 7, 38 (1967) 12)G.F. Freeguard, M. Karmarkar, J. Appl. Polym. Sci., 15, 1649 (1971) 13)K. Sardelis, H.J. Michels, G. Allen, J. Appl. Polym. Sci., 28, 3255 (1983) 14)G. Allen, M.J. Bowden, D.J. Blundell, F.G. Hutchinson, G.M. Jeffs, J. Vyroda, Polymer, 14, 597 (1973) 15)N. Shah, J. Mat. Sci., 23, 3623 (1988) 16)M.E. Fowler, H. Keskkule, D.R. Paul, Polymer, 28, 1703 (1987) 17)P. Laurienzo, M. Malinconico, E. Martuscelli, G. Ragosta, M.G. Volpe, Italian Patent N ~ 47946A89 18)P. Laurienzo, M. Malinconico, G. Ragosta, M.G. Volpe, Angew. Makromol. Chem. 170, 137 (1989) 19)G. Carbonara, P. Mormile, G. Abbate, U. Bemini, P. Maddalena, M. Malinconico, in "Physical Concepts of Materials for Novel Optoelectronic Device Applications I: Material Growth and Characterization",M. Razeghi ed., Proc. Soc. Photo-Opt. Instrum. Eng., ~

688 (1990)

468 20)U. Bemini, G. Carbonara, M. Malinconico, P. Mormile, P. Russo, M.G. Volpe, Appl. Opt., 31, 5794 (1992) 21)U. Bemini, P. Russo, M. Malinconico, E. Martuscelli, M.G. Volpe, P. Mormile, J. Mater. Sci., 28, 6399 (1993) 22)F. Coppola, R. Greco and G. Ragosta, J. Mater. Sci., 211775 (1986) 23)A.J. Kinloch and R.J. Young, "Fracture Behavior of Polymers", Applied Science, London, 1983 24)M.J. Doyle, J. Mater. Sci., 18, 687 (1983) 25)G. Abbate, U. Bemini, P. Maddalena, S. de Nicola, P. Mormile and G. Pierattini, Opt. Comm., 70(6), 502 (1989)

469

CHAPTER 9

POLYCARBONATE TOUGHENING BY ABS Roberto Greco

Institute of Research and Technology of Plastic Materials (IRTEMP) of Italian National Research Council (CNR) - Via Toiano, 6 - 80072 - Arco Felice (Napoli) Italy.

1. Introduction

Polycarbonate (PC) and Acrylonitrile-Butadiene-Styrene (ABS) blends are commercial products since many years; they have received a particular attention in patents and technical applications [1-3], such as in automotive industries. The reason of their success on the market is due to their excellent thermal, mechanical and impact performances. The two components offer, in fact, a good compensation of properties, as summarized in table 1, where positive and negative technological aspects of both PC and ABS are schematically illustrated. It is of a particular interest to the purpose of this paper to analyse the blend properties and particularly the impact performances, with respect to their simple and peculiar way of preparation. Their good mechanical and impact behaviours are obtained, in fact, by simple melt-mixings in common equipments, without addition of any specific additive for component compatibilization. This achievement can be considered rather an exception when compared with more complex techniques of toughening

used

for

other

incompatible

systems.

Suitable

mterfacial

compatibilizing agents must be generally added, in fact, to the mare components in order to reach the goal. Furthermore this is achieved by more or less complex

470 procedures, such as, for instance, reactive blendings. Several of these examples are illustrated elsewhere in this book. Therefore, since the behaviour of PC/ABS blends, is not yet clearly understood from a scientific point of view, it is worth to undertake a systematic investigation of these blends. In recent years, in fact, only a limited number of papers have been issued in literature, concerning some of their properties. As a matter of fact, PC/ABS alloys consist of four polymeric species, PC, polystyrene (PS), polyacrylonitrile (PAN) and polybutadiene (PB), the last three in form of copolymers, S-co-AN (SAN) and PAN grafted onto PB, all compounded in complex multiphase systems.

Table I Comparison of technological qualitative behaviours between PC and ABS technopolymers.

Behaviour

Positive

PC

"

ABS

high heat distortion T

economy

mechanical resistance

processability

low T toughness

impact strength

transparency

notch sensitivity

dimensional stability electric properties

Negative

processability notch sensitivity stress cracking chemical resistance

low heat distortion T

471 For a systematic study, necessary to highlight the intimate reasons of the blend performances, it is convenient to briefly describe chemical constitution, processing and properties of the single components. The successive step is that of describing the behaviour of PC/SAN blends before affording the more complex PC/ABS multicomponent systems. Finally the PC/ABS blend properties will be analysed and discussed. The acquired direct knowledge on PC/ABS blends could be useful for developing similar compatibilization techniques to be utilized for other systems.

2. Blend components

2.1 PC PC, consisting of linear thermoplastic polyesters of carbonic acid with aliphatic or aromatic dihydroxy compounds, can be represented by the general structure: O

II 0

C~O

A detailed description of all aliphatic and aromatic polycarbonates, a very large family of condensation polymers [4,5], is, however, well beyond the scope of this paper. Here only the Bisphenol-A Polycarbonate (PC-BPA), the most used polycarbonate for commercial blends with ABS is of interest, whose structure is the following:

L

o_c_o

PC-BPA will be simply indicated as PC in the remainder of this paper. It has a good thermal stability and the dry PC may be kept for hours at 310 ~ in the molten state, thermal degradation starting above 400 ~

472 Its injection moulding grades possess all average molecular weight (MW) generally comprised between 22000 and 32000. Tensile, impact and flexural strength increase with MW up to about 22000, beyond this value their improvement becomes much smaller, whereas the melt viscosity increases more sharply. Therefore a compromise is necessary to obtain a viscosity sufficiently high to get satisfactory mechanical properties but sufficiently low to impart flow characteristics suitable for filling complex moulds. PC is thermally and mechanically stable up to its Tg, lying around 150 ~ (its storage modulus, G', is only slightly dependent on temperature over this temperature interval). PC exhibits in solid state a toughness higher than other amorphous glassy polymers, such as polystyrene and polymethylmethacrylate. This has been attributed to a broad G" maximum, the so-called ~/-relaxation, occurring at about -100 ~ I61. A brittle-to-ductile transition occurs below room temperature. It depends on several variables, such as temperature, MW, loading rate, state of stress, thermal treatments (physical ageing and annealings), additives, external media, sample thickness and notch sharpness. PC, although exhibiting tough behaviour in stressstrain and in unnotched impact tests, is very sensitive to sharp notches and to specimen thickness.

2.2 SAN Styrene-acrylonitrile (SAN) copolymers are low-cost materials of increasing commercial importance, prepared from acrylonitrile I CH 2 = CHCN ] and styrene ICH2 =

CHC6Hs] monomers [7,8]. The monomers mixture is generally polymerised

according to the azeotropic ratio (76 wt % of styrene) in order to better control the copolymer composition.

473 SAN copolymers solve the problem of the bad processability of PAN, being easily moulded and shaped by conventional equipments. They are strong, rigid and transparent materials, with high dimensional stability, craze resistance and strong resistance to liquids, such as water, acid and basic aqueous solutions, detergents and bleaches as well as to solvents, such as oils, gasolines and kerosenes. They exhibit a better stress-cracking resistance than general-purpose PS in several environments, together with a good thermal stability, low creep behaviour, excellent tensile and flexural strength, surface hardness, great rigidity and good resistance to weather agents. Most of these characteristics are improved with increasing the AN amount in the copolymers, showing best behaviours in the range of 20-35% in weight of AN in SAN. Beyond this values a yellowing effect increases too, requiring suitable additives. The combination of the properties above illustrated has made SAN suitable for end-uses in several fields such as buildings, automotives, major and minor appliances in domestic equipments, packaging, home furnishings and others.

2.3 ABS

Acrylonitrile-butadiene-styrene (ABS) copolymers are a large family of thermoplastic

materials,

containing

an

elastomeric

component,

usually

polybutadiene (B) or a polybutadiene-based copolymer, in form of domains well dispersed in a thermoplastic matrix of SAN [9-11]. The SAN is graRed onto the elastomer in order to obtain a suitable dispersion and a reduction of the domain size. ABS can be produced by several techniques: a) mechanical blending of SAN and a SAN-co-B copolymer in common mixing equipments. b) polymerization: there are three commercial processes for ABS manufacturing:

474 b~) emulsion, involving a two step process: 1) production of an elastomeric substrate, made by PB, or Styrene-co-B (SBR) or AN-co-B (NBR) random copolymers; the S or AN amount must be less than 35 %, in order to keep the rubber glass transition temperature (Tg) contribution sufficiently low; the reaction is carried out in a batch reactor and a careful control is made on the particle size (in the range of 0.05-0.5 pan) and on the crosslinking degree, both affecting the gratting efficiency of the second stage of the process; 2) copolymerization of S and AN and simultaneous grafting reaction of SAN with the rubbery substrate formed in the first step. High rubber contents can be obtained by this technique ( up to 50 % ). These materials are often blended with SAN or other ABS materials to vary their initial rubber concentration and particle size distribution. Advantages are low temperatures and pressure used in the process and a wide range of products available. Disadvantages are high energy requirements. b2) suspension polymerization, involving a two-step process: 1) the rubber, dissolved in a styrene-acrylonitrile mixture together with a chain-transfer agent and an initiator, is charged to a prepolymerization reactor; the particle size (generally in the range of 0.5-5 lam) is partially controlled by adjusting MW and stirring intensity; the phase inversion sets the maximum rubber amount which can be incorporated since higher rubber contents increase too much theviscosity and consequently the average particle size; 2) the prepolymer is charged to a suspension reactor along with water, suspending agent, initiator and a chain-transfer agent. Water and energy consumption are lower, but the wastewater is more concentrated than for emulsion process.

b3) bulk polymerization, involving a continuous two-step process. A butadiene rubber is dissolved in a styrene-acrylonitrile mixture together with a transfer agent, an initiator and a diluent, for controlling the viscosity, in a

475 prepolymerization vessel. Discrete rubber particles including SAN and monomer are formed, whose size (0.5-10 ~tm) is controlled by a high shear stirring. The prepolymerized material is continuously charged to a polymerization reactor, where the rubber particles are cross-linked, retaining the shape previously acquired. The rubber percentage is limited up to 15-18 %, since, beyond this concentration, the viscosity becomes too high and the material processing too difficult. Advantages of this method are low energy requirements and low wastewater amounts and, hence, low costs of production. Disadvantages are high cost equipments, minor product flexibility, due strong limitations in processing highly viscous polymer melts, less complete conversion from monomer to polymer. This last effect requires, for most ABS, a devolatilization process, in order to freed them from residual monomer prior to compounding of the final product. All these methods of preparation yield a family of ABS materials with a large flexibility, depending on composition, MW, degree of graRmg, rubber particle size and morphology, allowing the tailoring of properties suitable to meet specific enduses. However, in spite of the different processing utilized, the morphology mainly consists, in all the cases, of a rubbery phase dispersed in a compatible way in a SAN matrix, due to the grafting of SAN onto the B-based rubber. ABS are engineering thermoplastics exhibiting good processability, excellent toughness and sufficient thermal stability. They have found application in many fields, such as appliances, building and construction, business machines, telephone, transportation, automotive industries, recreation, electronics and others.

3. Blends

3.1 PC/SAN blends A certain number of papers have analysed the properties of blends made by

476 SAN and PC [12-55] . One of the main aims, was to get information on their behaviour as a necessary step for affording the analysis of the more complex multicomponent PC/ABS systems. Keitz et al. found experimentally [12] that a number of properties reached a maximum value, for blend specimens having a AN percentage in SAN, ranging from 25 up to 27 % in weight: 1) lap shear adhesion of compression moulded laminated sheets of PC and SAN copolymers; 2) mechanical tensile modulus and elongation at break; 3) notched impact Izod strength; 4) inward shifts of the Tg of PC and SAN blend components with respect to the homopolymer values, detected by DSC and dynamic-mechanical tests. These findings were interpreted as due to PC and SAN adhesion, induced by a partial miscibility varying with SAN internal composition in the blends. A simple binary interaction model, already utilized for a copolymer, made of monomer 1 and 2, mixed with a pol3qner (3), was proposed for interpreting this behaviour

[12,13]. The overall interaction energy density B, as a function of three component binary interaction parameters, can be written as follows: B = B,3 ,J,', + B23 'b'2 - B,2 ,J,', 'b'2

(1)

where ~'~ and {D'2 a r e the monomer volume fractions in the copolymer. The first two terms, representing interaction parameters between polymer 3 with 1 and 2 monomers are linearly additive, as shown in Fig. 1. The third term, taking into account the intramolecular interaction parameter between the two monomers, is of a quadratic form. In the case of endothermic mixing, B~j in eq. 1 are all positive, then B, as a function of ~'~, tends to exhibit a minimmn ( dZB/d~}'l 2 - 2 B12 > 0 ).

477 Furthermore [ 14] for: B,2

>

(2)

(~mB-13 + -x/B23) 2

B becomes negative and therefore a miscibility window must exist in the composition range between +'~, and ~'~b, as shown in Fig. 1.

B Bls

B2s i

B12) 0 o

Figure 1.

1

Interaction parameter B versus ~b'~, for different B~e values,

greater than O, increasing in the arrow direction (after Keitz et al. [12]).

In our case SAN is the copolymer and PC the polymer: the maximum in the curve of lap-shear stress versus SAN composition could be an evidence of the existence of a miscibility window. It is interesting to note, for later discussion, that the above mentioned range of AN contents (25-27 wt %) is very close to the azeotropic composition of SAN, often used as matrix phase for ABS materials, as already mentioned above.

478 In Fig. 2 the lap-shear stress of adhesion, between PC and SAN sheets, is reported as a function of the AN percentage in SAN. !

I

I

I

I

I

10

20

30

40

50

60

_

T

(Pa)

0

70

AN (~)in SAN

Figure 2. Average lap-shear stress as a function of SAN internal composition (after Keitz et al. [12]).

Mendelson [15] studied the miscibility in blends of SAN copolymer with PC and several other polymeric species. PC and SAN were found to be phase separated but partially miscible in a broad concentration range of SAN (23-70% AN), as evidenced by Tg measurements. This result was in partial agreement with those of Keitz et al. [ 12] and of Locati et al. [ 16]. He assumed that a Tg linear model was applicable to each of the separated glassy phases (made by SAN-rich and PC-rich domains), as in the case of a homogeneous blend of two miscible polymers. He used the additivity model of a modified Gordon-Taylor equation [ 17], where the difference between the thermal expansion coefficients in liquid and solid states is taken as a constant for all polymers, as proposed by Tobolsky [18] :

479 Tglblend -- X

SAN

Tg

SAN

+ X PC Tg

PC

(3)

He calculated from the model that: a) less PC entered the SAN-rich phase than SAN the PC-rich phase; b) the SAN composition had only a small effect on the PC-SAN miscibility; that is, on the proportion of PC-rich and SAN-rich phases in the overall blend as well as on the distribution of PC and SAN between these two phases. This last finding was in contrast with the conclusions of Keitz et al. [12], for whom a maximum of miscibility existed in a restricted SAN composition range, as reported above. This could indicate that the application of the Tg model was not suitable for interpreting the data. TEM observations showed rather diffuse phase boundaries between PC and SAN domains. Both the findings, the conflicting results of the Tg model application as well as the TEM features, suggest a certain degree of interpenetration of PC and SAN domains across the interface. From mechanical tensile tests the amount of material deformed by crazing was found to decrease with increasing the PC content in the blend. Gregory et al. [19] analysed, by means of a torsion pendulum, the dynamicmechanical behaviour of multilayered composites, made of PC and SAN alternating laminates. Two series were made, consisting of 49 and 193 alternating layers of same overall thickness but of varying composition. The outer layers consisted of PC in all cases. A novel damping peak was observed in between those corresponding to the Tg of the two constituents. The presence of this peak was almost independent from a variety of parameters, such as molecular orientation, composition, thermal history, thermal cycling, number and thickness of the layers. The peak disappeared only when the planar structure of the layers was disrupted. Its origin was attributed to a particular temperature dependence of the viscoelastic parameters in the layer composite in the appropriate temperature regimes. The same authors investigated on the deformation behaviour of two series of similar coextruded multilayered composites [20] by macroscopic tensile tests,

480 performed at different strain rates. Optical microscopy was used to correlate the microscopic mechanisms observed in the two phases, with the modes of deformation observed in bulk. Three kinds of modes were observed (see Fig. 3) within a single bulk composition: 1) curve A~ :a brittle fracture at low strains, without yielding; 2) curve A2 : a ductile yielding followed by a rupture during the neck formation; 3) curve B : a ductile yielding with a stable neck formation followed by cold-drawing and rupture at high strains. The final mode of failure depended on the relative thickness of PC and SAN layers, as determined by composition and strain rate.

A2 B

F,

Figure 3. Typical modes of fracture of PCA7AN blends, observed by Gregory et al. [20].

Optical microscopy revealed craze initiations in the SAN layers which induced successively shear bands in the PC layers at the craze tips. This interaction between crazes and shear bands hampered the crack propagation delaying the rupture by a stress delocalization. Thermal, mechanical and impact properties of blends containing PC and two different types of SAN (containing 5.5 and 30% in weight of AN), obtained by

481 injection moulding were analysed by Skochdopole et al. [21]. Parameters, such as strength, modulus, heat distortion temperature under load, showed a linear dependence with blend composition. Others, such as elongation at break and dart impact energy dropped from PC down to SAN level at a percentage of about 40% of SAN. Izod impact toughness followed the same trend but the drop occurred at a much lower SAN content (only 10%). The Tg analysis evidenced only a limited solubility between the components, with the blend, containing SAN with a higher AN content (30%), the more soluble in PC. Berger et al. [22] studied the microdeformation at constant strain rate of casted thin films (0.4 lam) by optical microscopy and TEM. The average tensile strain for void formation, ev, was 0.13 for SAN and 0.23 for PC. Both PC and SAN underwent shear yielding, ev decreased by the PC addition to SAN up to 60% in the blend and the voids were initiated by crazes at the PC-SAN boundaries, ev decreased alter an annealing, performed at about 90~

and the shear yielding

mechanisms were suppressed, favouring crazes formation in SAN and in SAN-rich phases. A model, was proposed to interpret the phenomenon. Kim and Bums [23] examined the miscibility of PC and SAN by determining, on solution casted and extruded samples, Tg and ACp (Cp difference between rubbery and glassy state) of blends,. The Tg of PC decreased with increasing SAN content and viceversa that of SAN was enhanced with augmenting PC amount. ACp of both PC and SAN decreased with increasing amount of the second component. The overall behaviour was attributed to a partial dissolution of each component in the conjugate phase. From the experimental data of Tg and ACp the authors calculated the apparent weight fractions of the two components dissolved in the PC-rich and in SAN-rich phases. This was accomplished by the Couchman relationship [24,25] used to describe the Tg dependence on composition in miscible blends. It appeared that SAN dissolved more in PC-rich phase than PC in SAN-

482 rich phase, confirming the previous finding of Mendelson [15]. The polymerpolymer interaction parameter was calculated by the Flory-Huggings theory [26]. The extrudate swell exhibited a maximum and the viscosity a minimum at a weight blend ratio of 50/50, showing a positive and a negative deviation from a linear additivity rule respectively. A reasonable explanation was given for this double effect: when one of the fluids is the matrix, it flows mostly along the tube wall, dissipating more energy than the particles of the dispersed phase. These, on the other hand, being preferentiaY~,~transported along the centre of the tube, are elongated and store more elastic energy than the matrix. When the two phases are co-continuous a reciprocal lubricating effect and their deformation are maxima, giving rise to a minimum viscosity and a maximum die swell in the middle composition range. Guest and Daily [27] analysed the thermal behaviour of PC/SAN blends. The SAN was also dissolved in a suitable solvent and reprecipitated in order to remove from it all the low molecular weight species, that are generally present in styrenic

Table 2. Glass transition data for PC, SAN copolymers and (70/30) PC~SAN blends (after Guest et al. [27]). i

Sample

i

Tgpc,

ATpc

(of) PC

155

(oc) 5.0

SAN* Blend(70/30)*

152

8.0

SAN** Blend(70/30)**

155 * as received

\

TgSAN, ATsAN

6.0

108.5

6.75

113.5

7.0

114.5

5.5

114.5

5.5

** reprecipitated

483 polymers in amounts, ranging between 1 and 2 wt % [28]. A comparison was made between the Tg and the half-width of the glass transition, ATg, of PC and SAN, between "as received" and "reprecipitated" SAN and the corresponding blend values (see Table 2). As it is possible to note, only the "as received" materials show classical inward Tg shitts and ATg changes of blended PC and SAN with respect to the pure component values, whereas in the "reprecipitated" materials such variations are almost completely suppressed. Hence this phenomenon is clearly due to the migration of low molecular weight species, such as monomers and oligomers, contained in commercial SAN, towards PC domains. Also the viscosi~, ratio of the two components can be influenced by the presence of these plasticizers, determining in turn somewhat the blend morphology. In a successive work the same authors [29] showed the ability of dynamicmechanical spectra in revealing morphological features of PC/SAN blend obtained by different processing conditions, such as compression and injection moulding. The injection moulded samples exhibited, in fact, a Tg (G") peak, broader than that corresponding to compression moulded specimens, having the same composition. In other words, they showed somewhat an apparent greater degree of mixing. This effect was related to the mterfacial regions obtained by processing. A parallel analysis by SEM showed a coarse uniform structure throughout compression moulded samples. By comparison, injection moulded samples exhibited a similar morphology in the centre and a stratified, fine, laminar and oriented structure in the skin regions. This kind of morphology seems to resemble that of multilayer sheets studied by Gregory et al. [19, 20] , indicafng tlaat, also in this case, the layers continuity plays the major role. Compression and injection moulded specimens of PC/SAN blends were analysed by dynamic-mechanical torsional tests by Mclaughin [30] as well. The two kinds of specimens showed different morphological microstructures: a) for

484 compression moulding specimens, SAN spherical particles in a PC matrix at high PC/SAN ratios, or PC spherical particles in a SAN matrix at low PC/SAN ratios throughout the sample, viceversa; the onset of co-continuity occurred at a PC/SAN ratio, ranging between 50/50 and 60/40; b) for injection moulding specimens, spheroidal particles in the core and lamellae in the skin, at very high or very low PC/SAN ratios, and sheets all over the sample in an extended range between 50/50 and 80/20. An intermediate peak between the two Tg became distinguishable at high frequencies in plots of tan ~5versus temperature in the latter case. This finding was related to the co-continuity obtained in the injection moulded specimens, which could be idealised as alternating sheets, similar in structure and behaviour, to the previously mentioned multilayered composites analysed by Gregory et al. [ 19,20]. Huang et al. [31] , somewhat confirmed the previous suggestion [12] of an optimum compatibility between PC and SAN in blends with SAN, containing 25% in weight of AN, by means of theoretical calculations based on the solubility parameters of the two components. Quintens et al. [32,33] analysed the dependence of the viscoelastic properties from the morphology. They annealed, at a temperature higher than the Tg of PC (200~

two injection moulded samples of PC/SAN blends, of different

composition (70/30 and 60/40). A gradual coarsening of the microstructure was obtained with increasing the annealing time. As already described elsewhere [19,20,29,30], the storage modulus plateau increased in value in the temperature field ranging in between the two components Tg. Furthermore the coarsening induced a loss of ductility in tests at high deformations. These findings were attributed to the reduction of the PC/SAN interface, consequent to the induced morphological changes. The influence of SAN composition on tensile stress-strain behaviour and phase morphology of a 60/40 PC/SAN blends, containing SAN copolymers of different AN content (in the range 0-34 wt % of AN), was analysed by Quintens et

485 al. [34,35]. Maximum stress and elongation at fracture exhibited a maximum for blends with a SAN copolymer having 25% AN and the morphological dispersion was finest at this composition. A phase morphology coarsening, consequent to the thermal treatment described above, resulted in a loss of ductility but the maximum was not shifted, confirming previous fmdings [12,31]. Also in this case annealings at high temperatures destroyed SAN layers continuity produced during the injection moulding process. This occurred by the break up and coalescence of the elongated SAN particles, influencing the viscoelastic behaviour in between the Tg of the two components. The rate of the process was determined by the viscosity ratio of the two components at the temperature of annealing. The morphological trend, as a function of annealing time, could be monitored by the decreasing contribution of the SAN phase to the overall viscoelastic behaviour of the blend. Takahashi et al. investigated the rheological behaviour of several polymers (HDPE, PP, PS, Ny 66, PC, SAN) [36] and that of PC/SAN blends [37] at very high shear rates. For homopolymers the generalised flow curve, consisted of two non-Newtonian regions separated by a transition or by a Newtonian zone. In the first non-Newtonian zone the viscosity decrease was attributed to a lowering of entanglement density. In the second one the macromolecular chains underwent to a scission mechanism. Then, in spite of the molecular weight decrease, the chains were still able to reentangle with each other, giving rise to a sudden viscosity increase. Finally the viscosity decreased again at higher strain rates, due to a double effect: a) a new disentanglement process of the previously reentangled chains; b) a further chains snapping mechanism. For PC/SAN blends the trend was rather similar to that of the two components and a cylindrical multilayer model was proposed to fit the data. The effect of the high shear rate extrusion on the component compatibilization of a (70/30) PC/SAN blend was investigated by Takahashi et al. [38], as well. The apparent volume of the SAN particles dispersed in the PC matrix, decreased with

486 increasing shear rate, as shown by SEM observations,. Moreover repeated extrusions decreased this volume more and more and a shit~ of the Tg of PC was detected. This indicated that SAN could be partially dissolved into the PC matrix by means of an intense mechanical sheafing. Alternatively free radicals could be produced by the chains scissions. These could favourite the formation of interfacial agents between PC and SAN, helping the PC and SAN compatibilization. The emulsifying effect of the PC-g-SAN copolymers, formed in this way, would enhance the adhesion between PC and SAN, reduce the SAN domain sizes and create effective interlayer zones. Shah et al. [39] analysed ternary blends, formed by PC, one of two SAN (containing 13 % and 25 % of AN) and one of three aliphatic polyesters. Each of the three polyesters was capable to promote in binary blends a single Tg behaviour when mixed with SAN (25 % AN), SAN (13 % AN) or PC. All three of them were able to render miscible normally immiscible binary PC/SAN (25 % AN) or PC/SAN (13 % AN) blends. Each ternary blend showed, however, SAN-rich and PC-rich regions where immiscibility occurred. SAN (25% AN) was more easily solubilized in PC than SAN (13 % AN) by the best of the three polyesters. This indicates that the AN content plays an important role also in this phenomenon, confirming what previously found for binary PC/SAN blends [12,31,34,35]. Im et al. [40] presented a review on physicomechanical properties, such as tensile, impact and fatigue of PC/SAN microlayer composites. Particular attention was paid to irreversible deformations and damage mechanisms as well as to ductility improvements of the brittle components. Elastic moduli of compression and injection moulded PC/SAN blend specimens were measured at temperatures between the PC and SAN Tg by Arends [41]. In the intermediate blend composition range their values differed by about a factor of three, probably due to different modes of connectivity among the particles

487 of the nmlor component. The importance of phase continuity was evidenced for compression and moulding specimens by modulus data versus SAN composition. The data agreed quite well the behaviour predicted by an empirical relation based on the percolation approach in finite domains, devised from Monte Carlo simulations. The interaction between PC and SAN at a given AN content (about 25% by weight) was analysed again by Callaghan et al. [42] in regards to the molecular characteristics of PC and SAN. It was shown that the ~ o components become entirely miscible in blends when the molecular weight of each component is lower than about three thousand. The interaction energy was calculated by equation of state theories, developed by Flory-Huggms [26,43,44] and Sanchez-Lacombe [4550], allowing predictions of interfacial tensions.

3.5

I

'1

I

I

I

I

I

3.0 2.5

E

::L E~

J

2.0 1.5 1.0 0.5 t

I

I

I

I

I

I

5

10

15

20

25

30

35

40

AN (1~1 in SAN

Figure 4. Average diameter of dispersed SAN particles versus AN content in SAN in PC~SAN blends (after Callaghan et al. [42]).

488 The data are reported in Fig. 4 as particle average diameter, d~, versus AN wt % in SAN:

a minimum is observed at about 25 % of AN, consistent with

morphological observations made by Quintens et al. [34]. The interfacial tension appears to be directly proportional to the average SAN particle size in a PC matrix, provided that the components viscosity and the shear field are kept fixed, as shown also in some other case by Takeda et al. [51]. PC blends with commercial SAN purified from oligomeric species exhibited almost no measurable shifts in the Tg of either phase, as already found by Guest et al. [27], supporting their view of SAN oligomers partitioning between the two phases rather than the hypothesis of the partial miscibility made by several authors [12,15,21,23,31]. A series of SAN with different AN contents were selected in a way that their melt viscosity was maintained nearly constant and blended with a PC of a similar viscosity. In this conditions the morphology should be independent of rheological factors and should reflect only the interfacial tension. Mechanical properties of ternary PC/SAN/MBS (methacrylate-butadienestyrene copolymer) blends were analysed by Cheng et al. [52]. In particular the impact strength of a series of binary PC/SAN and ternary blends, containing SAN with different AN contents, as a function of this parameter, exhibited a maximum at about 25% of AN, as shown in Fig. 5. The blends were obtained by a simultaneous mixing procedure. Considerations of surface energy and component pair miscibility were successfully used to predict the MBS particles location. Most of the MBS impact modifier particles were found to be located at the PC/SAN interface, being trapped by surface forces. The mixing procedure as well as mechanical and impact deformation mechanisms were important, however, in determining this location in some cases.

489 500

- .9 E

I

I

I

I

I

I

400

--3

..c

300

C"

k--

"-'

200

E

100 ( _

PC/SAN

( 0

0

j

J

J

I,

J

5

10

15

2O

25

I

30

35

AN (~)in SAN Figure 5. Impact strength of SAN-based materials as a funcaon of AN percentage m SAN for PC/SAN/MBS and PC/SAN blends, as indicated (after Cheng et aL [52]).

The interfacial tension between PC and a series of SAN copolymers having AN contents from 0 to 40 % in weight were determined by measuring capillary thread ruptures at 200~ by Watkins et al. [53]. The interfacial tension plotted versus AN % exhibited a minimum at about 15 wt % of AN, a value lower of that previously reported [12,31,34,35,39,42,52] for the optimum value of miscibility between PC and SAN (about 25 wt %). When the interfacial tension was at a minimum and PC and SAN viscosities were similar, very thin threads could last for several minutes before rupture. This could be the reason of the formation of filaments and/or of stratified lamellae in the skin of injection moulded samples under the action of both high, shear and extensional, flow fields.

490 Inward Tg shifts of PC and SAN were detected on PC/SAN blends and the Flory-Huggins parameters were calculated by Kolarik et al. [54]. The compositions of the conjugate phases were calculated by the Fox's equation. The apparent solubility of SAN in PC or PC in SAN increased with decreasing PC or SAN volume fraction in blends respectively. These dependencies were tentatively ascribed to the presence of intermixed zones existing at the boundaries between PC and SAN domains. The yield stress followed a classical rule of mixing, indicating that the phase adhesion was sufficiently good to provide effective stress transfer between the phase domains during yielding and cold drawing. Janarthanan et al. [55] confirmed the existence of a maximum in a plot of interfacial fracture toughness versus AN % in SAN for PC/SAN blends at about 24 wt.% of AN in SAN, as previously found by Keitz et al. [12].

20 16

G

12 r

{J/m z)

AN '

I

'

2

I

4

'

I

6

.

I

I

8

Wt ~ of benzonitrile

Figure 6. Fracture toughness of PC~SAN interface, Gc, versus benzonitrile content for two different values of AN % (after Janarthanan et al. [55]).

491 They showed in addition, by using benzonitrile as a model oligomer, that low MW SAN species preferentially migrate toward the PC/SAN interface. This increases the PC/SAN interphase thickness reducing the entanglements between PC and SAN macromolecules. In fig. 6 the fracture toughness of adhesion between PC and SAN is reported against the amount of benzonitrile for two blends containing SAN with a different AN wt %. Both the curves undergo to a monotonic decrease, starting from different initial values depending on the AN % in SAN. All the literature results, above described, can be briefly summarized as follows: a) PC and SAN are completely immiscible in blends, both in melt [23] as well as in solid state [ 12,41 ]. b) Their blends show inward shifts of blend Tg values with respect to homopolymers ones. These Tg shifts have been often interpreted by several authors as due to partial miscibility [12,15,21,23,31,54] of the components. The real cause is due, indeed, to migrations of SAN low M.W. species towards PC domains during the melt-mixing [27,42,55]. c) There is a narrow interval of SAN composition (around 25 wt % of AN in SAN), where a variety of overall properties and particularly, the adhesion between PC and SAN, exhibit maxima [12,31,34,35,39,42,52,53,55]. This behaviour has been explained by a binary interaction model [12,13]. The interfacial tension between PC and SAN shows a minimum in the same AN composition range in SAN [12,42,52,53,55]. d) Low M.W. SAN species tend to migrate from SAN phase towards PC and enrich the PC/SAN interface, diluting the macromolecular concentration in the interzones. The net result is a decrease of the number of molecular entanglements between PC and SAN in the interzones at the PC/SAN boundaries. This lowered interconnectivity decreases, in turn, the adhesion between PC and SAN phases I55].

492 e) Different blend processings (compression and injection moulding), can give diverse morphological features, which can be monitored by dynamic-mechanical tests, as shown by several authors [19,20,29,30,32,33,41]. The storage modulus (G') versus T shows measurable changes (a new peak or a shoulder) in between the Tg of the two components. This suggests that PC and SAN changes of Tg in blends depended not only on the thermodynamics of the interface but also on the final morphology of the systems (changing, for instance, the surface to volume ratio). Therefore any external change, such as thermal treatments, yielding variations in the morphological characteristics of the systems, can be easily detected by this technique.

3.2 PC/ABS blends

The behaviour of a 50/50 commercial blend by nuclear magnetic resonance (NMR) in a range of temperature from 100 ~ up to 500 ~

by SEM and TEM

was analysed by Stefan et al. [56]. The transitions associated with PC, SAN, and PB were monitored. It was found that, at temperatures close to the Tg of PC, the chain motions of the zones surrounding the PC were strongly influenced by the PC itself. Therefore the PC apparent volume (interaction zone) increased of about 50% in comparison with its actual value. At temperatures, close to the Tg of PB, the rubber affected the surrounding regions even at a higher extent, so that the PB apparent volume increased twice its actual volume. The larger interaction of the rubber was attributed to the diverse morphology (small round particles), as well as to the higher flexibility of the PB chains. The improved blend impact performances with respect to pure PC was related to the apparent volume of the rubber phase. These findings suggests the existence if interacting regions between PC and ABS phases. The melt viscosity of PC/ABS blends as a function of composition, measured by Dobrescu et al. [57], is shown in Fig. 7. The viscosity goes through a minimum,

493 both at low (y - 1 sec -1) and high shear rate (y = 103 sec-1). The depth of the minimum is the larger the lower the shear rate. The composition value at which the minimum occurs seems to depend on the shear rate value. The trend of the curve is due the PC/ABS immiscibility in the melt. This effect induces a reciprocal lubricating effect of the two components during the mixing, lowering the overall internal friction of the material.

8000

k,~

"

I

'

I

'

!

'

I

'

I

'

'/

'

I

"'

i

''

I

' ........

CtJ

a.

0~

800

0 o

80

0

10

20

30

40

50

60

70

80

90

100

Wt ~; ABS or PMMA

Figure 7. Viscosity as a function of ABS or PMMA percentage for PC/ABS (full circles) and for PC/PMMA (empty circles) at two different shear rates as indicated (after Dobrescu et aL [57]).

In the same figure the curves relative to PC/PMMA blends are reported by comparison as well. They present no minima but a smooth decreasing trend,

494 particularly at high shear rates. The different behaviour of PC/PMMA blends with respect to PC/ABS ones depends on the PC and PMMA miscibility. Kim and Bums [23] found that the viscosity of PC/ABS blends, measured in a capillary rheometer, went through a minimum at a weight blend ratio of 50/50, as found by Dobrescu et al. [57] (see Fig. 7). In Fig. 8, the extrudate swell ratio is reported as a function of SAN, ABS or Kodar

contents.

The

last

one

is

a

copolyester

formed

from

1.4-

cyclohexanedimethanol and a mixture of terephtalic and isophtalic acids, which is known to be compatible with PC.

1.8 9"" - ' ~ " "" "-=.,.,~j~ P C / ABS 0

1.6

L_

~ ~5

1.4

1.2

1

0

0.2

0.4

0.6

0.8

1

Wt ~ SAN, ABS or Kodar

Figure 8. Die swell ratio as a function of SAN, ABS or Kodar percentage (after Kim et al. [23]).

495 The comparison between the three kinds of blends shows a very pronounced maximum for both PC/SAN and PC/ABS blends, at about 50 % of PC content. The third component exhibit, instead, a very fiat curve. These findings clearly indicates immiscibility in melt between PC and SAN or ABS and compatibility between PC and Kodar. The reasons of this behaviour have been already illustrated above for PC/SAN blends and they do not change in the case of PC/ABS ones. Both the Tg of PC and ABS of PC/ABS blends were measured by DSC. That of PC decreased almost linearly with increasing the ABS content. That of ABS, on the other hand, increased linearly with increasing the PC content (in an analogous manner as described for PC/SAN blends).

150

"i'

I

'

I

'

I

'

140

(}.. ~"

ABS " -9 0

130

i-

120

_.._~_.~ABS

110

......... 100

II~|

0

~

I

10

~

~ I

20

,,

i

30

S ,

i

40

,

A i

50

9

N i

60

,

i

70

,

i

80

9

i

90

100

PC (%}

Figure 9. Comparison between Tg of PC (circles) and of SAN or ABS (squares)" for PC/SAN blends (filled symbols) and for PC/ABS blends (empty symbols) (after Kim et al. [23]).

496 The ATg value of PC was greater than that of SAN at equal composition values, whereas the ATg of SAN and ABS followed a parallel trend, as shown in Fig. 9. This effect was attributed to the polybutadiene chains acting as an additional plasticizer for PC. The polymer-polymer interaction parameter was calculated also for PC/ABS blends by the Flory-Huggins theory, showing very close values to those calculated for PS/SAN blends. Cooney [58] tested the photostability of films of a commercial PC/ABS blend, found to be rather sensitive to UV and visible radiations. Particularly the polybutadiene on the specimen surface was easily oxidised and crosslinked. This effect embrittled the surface, causing superficial cracks on bending, which lowered the material impact strength. Suarez et al. [59] found that, for extruded sheets as well as for injection moulded bars, the modulus and the tensile yield strength were nearly additive whereas the elongation at break showed a minimum as a function of the composition. The notched Izod impact strength was much lower for ABS than for PC. It levelled off at the ABS value up to 50% of PC content and then increased almost linearly up to the PC value. Weber et al. [60] studied the morphology of PC/ABS blends in relation to the PC/SAN composition (keeping constant the amount of the SAN-g-B copolymer) and to the melt processing temperature. The Vicat temperature slightly depended on the melt temperature and increased with increasing the PC content. The morphology and the impact resistance were strongly influenced not only by the blend composition but also by the melt temperature. An enhancement of latter parameter caused, indeed, a demixing of PC and SAN and a worsening of the impact performances. However, when PC was the matrix this effect became negligible.

497 The influence of reprocessing PC/ABS blends on their physical properties was studied by Eguiazabal et al. [61]. Two processing cycles affected only slightly the properties. For a higher cycle number a change of the rubbery phase (cross-linking and oxidation) was observed. Therefore density, MFI, stress and elongation at break showed drastic variations after the first two cycles, whereas the small deformation properties were almost unaltered by reprocessmg. The toughening mechanisms of these blends were analysed by Ishikawa et al. [62] by three point bending of round notched bars. The addition of small ABS amounts (2%) decreased the PC toughness whereas contents from 5% up to 20% yielded larger and larger improvements. A stress whitening effect was attributed to formation of voids due to the fracture of the interphase existing between PC and ABS. Radusch et al. [63], studied dynamic-mechanical and dielectric properties, suggesting the existence of a partial compatibility between PC and ABS in the boundary layers of the order of magnitude of 1 nm. The macroscopic properties were, therefore, determined by the interactions among the phases in this layers. An optimal concentration for toughened behaviour was about 70 % PC. Lee et al. analysed the solid-state morphology of an injection moulded PC/ABS (90/10) blend [64]. A bead-and-string structure was observed in the skin regions of the plaques, whereas in the middle an isotropic ABS phase was dispersed in the PC matrix. In a subsequent paper [65] the analysis was extended to the entire blend composition. Three composition ranges were identified: a) a PCrich blend, where the situation was that described above; b) a mid range (between PC/ABS 70/30 and 60/40), where the previous structure evolved to a coalesced stratified morphology at the edges and to a coarse dispersed ABS phase in the centre, with some regions of co-continuity with PC; c) an ABS-rich blend with dispersed PC domains. Qualitatively the above results were explained as due to the

498 melt flow pattems during the mould filling and to relaxations and coalescence acting during the cooling of the materials prior to the complete solidification. Chun et al. [66] found a synergistic effect in notched Izod impact strength extended in a broad composition range (PC/ABS from 80/20 up to 10/90), wider than in other authors' findings. This was attributed the PC/SAN miscibility, to a suitable ABS composition (low B content) and to an improved mixing device. Triphenyl phosphate and a brominate phosphate were compared as flame retardant additives for PC/ABS blends by Green [67]. The latter was shown to be more efficient as flame retardant and to give higher heat distortion temperatures than the former. Real time small-angle X-ray scattering was performed by Bubeck et al. [68] on a series of rubber-modified thermoplastics in order to investigate the modes of deformation in tensile impact tests in such materials. In particular for PC/ABS blends the predominant mechanisms were shear yielding in the PC and associated rubber particle cavitation in the ABS. The cavitation mode seemed to provide a direct relationship between rubber content and impact toughness. The mechanism did not change whether the tensile impact direction was perpendicular or parallel to that of the injection-moulding. At last crazing, though precursor of the final fracture, occurred after the prevailing non-crazing mechanisms, contributing only for a few percent to the total plastic deformation. Suyun [69] found that the fracture morphology of PC/ABS blends with high notched impact strength was a synergistic combination of the rupture characteristics of PC (smooth surfaces and striation with branches) and ABS (microcavities and parabolic markings). The fracture morphology of blends with low impact strengths was very different from that of the two components. Both fracture morphology and impact strength were dependent on PC and ABS characteristics as well as on blend composition.

499 The viscosity-composition relationship of PC/ABS systems was analysed by Kumar et al. [70] by using the Lecyar model. This model is a simplified version of the more accurate Mc Allister's one. The experimental data showed a strong negative deviation from the linear trend with respect to the homopolymer values, as found by other authors [23,57,58]. A good fit of the data with the model was obtained. An accurate fractographic stud), was performed by Lee et al. [71] on PC/ABS single notched specimens obtained by injection moulding and fractured in tensile mode in an Instron Machine at a relatively low strain rate (48 mm/sec). Stress whitening was observed on the surfaces of ABS and of all the blends but PC, indicating voiding formation during the deformation and fracture. Plane stress flow lines were observed for PC, PC/ABS 90/10 and 70/30, accompanied by a lateral contraction of the overall sample. Also the 50/50 blend exhibited some features of plato stress at the specimen's edges. In the centre, however, a valley on one surface and a corresponding ridge on the conjugate surface evidenced a plain strata region. The resulting mechanism was therefore a sort of mixed fracture mode. No suckingin (overall contraction) was observed on 30/70 and 10/90 PC/ABS blends. A characteristic feature (called herringbone or chevron fracture) was however still visible in the centre of the specimen with narrow shear lips. Also this kind of failure mechanism was considered to be of a ductile nature. The ABS surface showed no sucking and no shear lips at the edge, indicating a macroscopic condition of plain strain. From the above observations a progression in the fracture mode with varying the composition was described. There was, in fact, a gradual change from a shear fracture of PC under plain stress conditions, to a craze failure of ABS, under plain strain conditions. The mare influence of the ABS addition to a PC matrix was the cavitational mechanism of the rubber particles. The plain strain observed in the central region of the specimens in intermediate compositions, was due to the increase of the ABS content, while the shear lips at the edges,

500 characteristic of the plain stress conditions became narrower, completely disappearing for pure ABS. An S-shaped curve was observed between the ductileto-brittle transition temperature

and the

composition.

The

most

ductile

compositions were PC/ABS 70/30 and 60/40 whereas the most brittle were 30/70 and 10/90. In a subsequent paper Lee et al. [72], performed a fractographic and a morphological analysis on an injection moulded specimen of PC/ABS composition 70/30. Fracture occurring perpendicularly to the injection direction yielded a herringbone feature in the surface. The one occurring parallel to it yielded an inverse herringbone feature. The blend toughness decreased in the second case. The herringbone was determined by interactions of the main crack with secondary cracks started along the centre line. The inverse herringbone had the same origin but the secondary cracks were initiated near the edges. Such differences in behaviour between perpendicular and parallel directions were attributed to the PC orientation induced by the shear flow during the processing. Vibration welding was applied by Stokes et al. [73] to ABS blended with other thermoplastics, among which PC. It was found that the weld strength of ABS, coupled with PC, was 92% of the value obtained by ABS coupled with itself. Moreover the fracture in tensile mode, performed in direction perpendicular to the weld surface, let~ on the PC halves several ABS islands. Both the findings appeared to be a further evidence of the good adhesion existing between the two resins. A systematic work on PC/ABS blends was initiated a few years ago by Greco et al., whose preliminary data have been published in several papers [74-77]. In a first work [74] two complementary techniques, suitable to selectively etch one of the two components, were presented. The aim was to get a phase contrast suitable for accurate

scanning electron microscopy

(SEM)

observations,

501 particularly in the composition ranges were one of the component is the matrix and the other is dispersed in it. In this case a cross-check is a convenient method. A unique etching technique would give, in fact, good information only for compositions where the phase to be etched is the dispersed one. In the opposite case, the result would be rather uncertain due to loss of the remaining dispersed particles from the etched substrate (matrix).

Figure 10. SEM micrographs of smoothed surfaces ~f a 70/30 (PC/ABS) blend: a) PC matrix etched by a NaOH aqueous solution; b) ABS dispersed particles etched by an acid aqueous solution (after Greco et al. [74])

502 In figs. 10 morphological features evidenced by two etching techniques are shown for a 70/30 (PC/ABS) blend, first etched and then coated with an Au/Pd alloy fihn, for scanning electron microscopy observations. In Fig. 10a the PC matrix has been completely etched by a PC hydrolysis, induced by a NaOH aqueous solution (30% w/v of NaOH), leaving on the surface ABS dispersed particles. Their surfaces are not completely smooth, as one would expect from immiscible components, but several small holes are present on them, due to complementary PC etched protrusions. This indicates a very extended interfacial contact area existing between PC and ABS phases. In Fig. 10b the complementary etching has le~ unaltered the PC matrix and therefore, the numerous holes, visible on the surface, represent the sites where the

Figure 11. SEM micrographs of smoothed surfaces of a 40/60 (PC/ABS), cut along two orthogonal directions, etched by an acid solution (after Greco et al. [74]).

503 ABS particles were sitting before their etching out from the matrix, by a strong oxidant acid solution. In Fig. 11 SEM micrograph of two sections, perpendicular to each other, of the same 40/60 (PC/ABS) blend are shown: the etching, made by the acid solution, reveals the shape of cylindrical PC domains in the ABS matrix. From a complete set of micrographs, encompassing all the blend composition, it was possible to establish at which composition a phase inversion between PC and ABS occurred. In a second paper [75] a brief literature review on PC/SAN and PC/ABS properties was presented, of which this paper represents a more extended version. 100

.

.

,

.

,

,

7O 2~ 'I--

50 4O 30

I-

20

]--

"~"~"~'..._

"-/

16

o"'~-~-"'-"-.----.~

_

93 2

10

"-64 0

20

40

60

80

100

ABS (%1

Figure 12. Ratio of mixing torque, T, to roller speed of rotation, n, (T/n), as a function of PC % and n values in PC/ABS(M) blends' (after Greco et al. [761).

The ratio of a Brabender-like torque (T) to roller speed (n), as a function of

504 ABS content, decreases with the ABS addition from the PC value at all n (at low values the curves exhibit a slight decrease with a series of intermediate peaks). For high roller speeds (n - 32-64 rpm) the curves, alter an initial sharp decrease, levels off on the final ABS value at a blend composition of about 40 % of ABS.

160

'

I

'

l

'

I

'

I

i

PC

140

1:3

~O o w

I-120 B,,....~

[] rl

ABS

100

0

20

---O.~

' - " ~ - ' - "---fi--- - .~

40

60

[]

80

~---,

100

ABS (%) Figure 13. Tg of PC and ABSfrom different sources: full circles, (Kim et al. [23]); empty squares, (Greco et al. [76]).

Processability, thermal, mechanical and impact behaviour were analysed [76] successively on the same PC/ABS(M) blend utilized in the first paper [74]. Improved PC processability by ABS addition was confirmed, as shown in Fig. 9. Inward Tg shifts, as found by other authors, both for PC/SAN [12,15,16, 21,23,27,42,54] and PC/ABS [23,60] systems, were detected by DSC. The data (squares) are shown in Fig. 13, compared with those of Kim et al. [23] (circles).

505 The PC Tg lie very close; the ABS ones follow a parallel trend. This behaviour can be due to a similar ABS internal concentration of the two types of blends. Stress-strain curves of PC, ABS and PC/ABS blends, relative to tensile tests made at R.T. on unnotched specimens, at a low deformation rate (0.1 mm/min), are shown in Fig. 14. Going from pure PC and adding more and more ABS the yield peaks lower and broaden down to a PC content of about 50 %; below this value the curve shapes resemble those of pure ABS, indicating that a phase inversion occurs, as confirmed by morphological observations as well [74]. The yield peaks and the elongation at break lowering, with decreasing the PC content in the blend, are evidences of a decreased blend ductility. 60

50

0

PC....,

. . . . . .

PC-8O

40 ~_~0-20 ~: "--" IO

30

\

~ AB S

PC-lO

PC-50 PC-40

2O 10 0

=

0

Figure 14.

I

10

i

I

20

,

I

30

=

I

40

=

I

50

Stress-strain curves of PC, ABS and PC/ABS(M) blends as

indicated (after Greco et al. [76]).

The application of the Kemer's model [78-80] to the experimental data of the Young modulus showed a very good adhesion between PC and ABS domains all over the composition range. The experimental points (empty circles) lie, in

506 fact, along the lines of perfect adhesion of the Kemer's model, as shown in Fig. 15,

1.4

/

'

'

'

'

/

L 1.2

'

/

Perfect

/

o

'

'

'

'

/ /

Adhesion

_1

o

1.0 ---~ LU

o Adhesion 0.8

0.6

0.4

0

0.2

0.4 ABS

volume

0.6

0.8

1

fraction

Figure 15. Curves of the Kerner's models for perfect and no adhesion between the phases, compared with experimental points (empty circles) (after Greco et al. [76]).

where the model curves of no (very poor) adhesion are reported as well. A synergistic effect was observed for Charpy impact strength and for maximum impact stress at about 25 wt. % of ABS, as shown in Fig. 14. A series of transitions from plain-strain to plain-stress conditions were observed at the different compositions, as clearly shown in Fig. 16. Starting from pure PC a brittle fracture (B) was observed for PC amounts ranging from 100 up to 80 wt %,

507 followed by a very ductile zone (D) between 80 and 50 % of PC (a strong synergistic effect is represented by two peaks in both E and ~ at about 75% of PC content).

12

,

11

- B

10

-

i

,

i

,

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i, !

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: ! !

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SD

9 8 OL.

7

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-

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0

~

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.,--"J~'~----"~

ILl

!

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20

~

I

t

40

I

60

t

I!

80

100

ABS(M)

l~gure 16. Max. stress (left hand-side axis) and energy to crack initiation (right hand-side axis) versus PC content. Plot impact behaviour zones: D, ductile; SD, semiductile; B, brittle (after Greco et al. [76]).

With a further increase of the ABS content (PC content between 50 and 20%) again the behaviour becomes brittle (B) and finally the ABS semiductilc (SD) fracture is approached. This alternation of impact mechanisms is directly detcctable by a parallel fractographic analysis. In Figs. 17 a few micrographs show fracture surfaces of specimens representative of the four composition zones, exhibiting an alternation of impact behaviour. The first one (a) shows classical features of a PC surface fracture, with

508

Figure 17. SF)~/Imicrographs ~ (a), PC; (b), 75/25 and (c), 50/50, PC/ABS blends ; (d) ABS (after Greco el al. [76]).

509 a relatively slow crack initiation region, where ribs are visible, and a successive zone of very high propagation rate, resulting in a smooth region. In this case plainstrain conditions can be judged from the inspection of the overall specimen shape as well, exhibiting no lateral contraction. Next surface (b), relative to a blend, containing 75 % of PC, corresponds to the g and E peaks (zone D) in Fig. 16: patterns of flow lines and a marked lateral contraction of the specimen are a clear evidence of plain-stress conditions. Next surface (c), of a rough appearance, represents a blend, made of 50 % of PC, where again the breakage occurs under plain-strain conditions. The last one (d) is relative to the semiductile behaviour of pure ABS, where slight flow lines reveal that a certain amount of material deformation occurs all over the surface prior to the crack opening. A comparison of the above illustrated results with those obtained using the same PC blended with a second ABS(B) in the PC/ABS blends was presented in the last paper of the series [77].

Table 2. ,Some characteristics of three ABS used in PC/ABS blend~ [77, 81]. i

TRADE

CODE

AN %

B%

S%

NAME Sinkral

MFI (ASTM D-1238)

B32

B

27

22

51

4

Sinkral M 122

M

23

14

63

18

Sinkral

A

27

11.5

61.5

26

A12

In the present paper processability, thermal and impact behaviour of a third ABS(A) has been added to the comparison. Extended trade name, code, internal composition, mclt flow index (MFI) calculated according ASTM D-1238, of the

510 three commercial ABS, manufactured by Enichem Inc., with the common tradename of Sinkral, are reported in Table 3. The main differences of the used ABS consists in the B content, decreasing from B (22%) to M (14%) to A (11.5 %), as well as in the MFI, increasing in the same order. The two ABS, M and A, are rather similar in composition being slightly different only in the AN content and in the MFI. The values of all three, however, lie very close to the AN range of maximum adhesion between PC and SAN in PC/SAN blends [12,31,34,35,39,52,53, 55].

500

I

'

'

'

I

'

'

'

'I

'

'

'

I

,

w

'

'I

_

_

Q_

c

50 tI--

ABS{A)

BS(B) -

;

{M)

I

I

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,

I

1

,

I

I

I

i

lO

I

I

I

I

I

,

,

-

i

I

10o

n (s - 1 )

Figure 18. T/n as a function of n for PC/ABS blends containing three ABS of different internal composition, M, B and A, as indicated in the figure and in Table 3 (after Greco et al. [77, 81]).

A more complete rheological characterization is provided in Fig. 18, where

511 the ratio of the Brabender torque per unit volume to the roller rotation speed, T/n, with the dimensions of an apparent viscosity (MPa-s), is reported as a function of n for PC and the three examined ABS. The PC has a less pseudoplastic behaviour than the three ABS, of which the ABS(B), having a higher B content, exhibits a higher internal friction (T/n has the same dimensions of an apparent viscosity, even though the complex pattem of the mixing process is completely different from those of common viscometers) than the other two ABS, during the treatment in a Brabender-like mixer. Moreover its curve crosses the PC one at an n value of 16 rpm. Therefore, at the rotation speed at which the blends were prepared (32 rpm), its "viscosity" is

45 40

=

2

rpm

35

ao ~

I---

-

25 20 -

15 10

-

M i

I

20

i

I

I

40

I

60

i

I

80

n

100

ABS (%)

Figure 19. Ratio of torque, T, to roller speed of rotation, n, (T/n), at n value of 32 rpm, as a function of PC %, for blends containing different ABS; codes as indicated in table 3 (after Greco et al. [77, 81]).

512 lower than that of PC. The other two ABS show very close values with respect to each other but much lower than the PC ones. These rheological features can have great influence in the processability of the blends and, of course on their final morphology and properties. As a reference to real conditions of the blend mixing, in Fig. 19, the curves of the same viscometric parameter, T/n, have been reported as a function of blend composition. As it is possible to see, all three ABS reduce the PC internal friction, greatly improving the PC processability when PC is the matrix of the system. The best effect is provided by the ABS(B), the most viscous among the three ABS probably because of its higher B content. The greater B amount would enhance the reciprocal lubricating effect necessary to lower the internal friction of the system. With increasing the ABS content a phase inversion occurs and, therefore, the properties of the ABS matrix become predominant for all the blends. For PC/ABS(M) and PC/ABS(A) the internal friction is very low at high ABS contents compared with the PC one at 32 rpm (mixing roller speed) and T/n levels off on the ABS value. Only the PC/ABS(B) exhibits a very pronounced minimum, since, when ABS(B) is matrix, T/n must gradually come up to the high ABS value, due to its large internal friction, comparable with that of PC. The thermal behaviour of the three types of blends is reported in Fig. 20, where the Tg of PC and ABS are plotted as a function of the blend composition. The most relevant feature if the higher ATg of PC, between the blend PC/ABS(B) and the PC homopolymer, with increasing the amount of ABS, in comparison with the two other types of blends, M and A. This effect somewhat confirms the marked contribution to Tg shitts provided by the B, previously illustrated in Fig. 7 [23], where the ATg of PC was larger for PC/ABS than for PC/SAN blends (containing no B).

513 A migration of low MW species of B (m addition to those of SAN species [27]) from the ABS domains towards those of PC along the PC/ABS boundaries was invoked by the authors when the SAN component was substituted by the ABS one. Also in our case it seems that the higher the B amount in the ABS, as it is for ABS(B) in regards to ABS(M) and ABS(A), the larger the contribution given by this component to the PC plasticization, in the PC/ABS interzones. 155

,

,

,

,

,

,

;

,

,

,

,

,

,

,"

,

""

145

M

0

135

A

Q

I--

125

A M

115

o

B 105

,

0

i

---L.___.____

9

0

9

!

20

i

,

I

I

40

,

I

60

Ii

,

i

!

80

,

,,

i

100

ABS % Figure 20. PC and ABS Tg versus ABS % for PC/ABS blends, containing different ABS, codes as indicated in table 3 (after Greco et al. [77, 81]).

The impact performance (as energy to crack initiation, E as a function of blend composition) of the three blends, tested in flexural Charpy mode on sharply notched specimens, are compared in Fig. 21. The results are similar for PC/ABS(M) and for PC/ABS(B), with some significant differences: a) the synergistic peak, present in both materials, has a different shape (sharp and tall for the former, broad and low for the latter); b) the

514 composition at which the first brittle-to-ductile transition occurs (see Fig. 14) is lower for ABS(B) than for ABS(M): only 10% of ABS is sufficient to toughen the PC in the first case, whereas in the second case it is necessary add at least 20 % ABS to reach similar results. This seems to be due the higher B content of the former. The third blend PC/ABS(A) shows no synergistic effect in all the blend composition, probably because of a different morphology. 10

8

&-"

6

uJ

4

2

0

0

20

40 ABS

60

80

100

1~}

Figure 21. Impact Charpy strength as a function of ABS content for different ABS, codes as indicated in Table 3 (after Greco et al. [77, 81]).

A more complete and accurate analysis on the effect of the ABS intemal composition will be presented in a forthcoming paper [81]. Other authors have analysed the influgnce of the ABS composition on PC/ABS blends, as well, even though in a non systematic way.

515 Kurauchi et al. [82] analysed two blends, of a commercial PC with two commercial ABS, having a different (AN/B/S) composition: a) 22/29/49 ; b) 20/37/43. Only the blend made with the first ABS (a) showed synergistic effects in tensile stress-strain parameters (energy absorption and elongation at break) and in unnotched impact Charpy strength, exhibiting maxima at high PC contents, as a function of composition. In the other blend, with ABS (b), having a greater rubber amount than (a), the corresponding parameters showed, instead, an almost linear trend at high PC contents and minima at low PC amounts. This effect could be due to the different B amounts of the two ABS(a, 29% and b, 37%). In the second case, in fact, the B could encapsulate most of the SAN, reducing the PC/SAN surface of contact. In addition, as a secondary effect, the AN content in b) is slightly lower than that in a), reducing in the first case the PC affinity with SAN. Morbitzer et al. [83] analysed thermal, dynamic-mechanical, stress-strain and notched impact behaviours of two PC/ABS blends, obtained according two different procedures: I) SAN and a SAN-g-B copolymer were preblended in a 60/40 ratio and then various amounts of PC were added to this blend; II) PC and SAN were preblended in a 50/50 ratio and then different amounts of a SAN-g-S copolymer were added to this blend. In both cases the Tg of PC, SAN and B, clearly detected by thermal and by dynamic-mechanical techniques, changed their values with varying the blend composition. The first two Tg varied as reported by other authors for PC/SAN blends [7,10,16,18,23,37,41 ]. The Tg of B decreased to lower temperatures with increasing the PC amount (decreasing the ABS and consequently the B amount). This effect was attributed to the thermal stresses caused during the blend cooling by the different thermal expansion coefficients of the grafted rubber particles and the surrounding SAN matrix [84]. The stress fields in the matrix, around the rubber particles, tended to overlap in dependence of the rubber content and of the final morphology. Since

516 these stresses were lower than the interfacial forces, the rubber particles underwent to a negative hydrostatic pressure. This induced an increase in free volume and a consequent decrease in Tg. The stress and the elongation at break, measured at room temperature, as well as the impact strength, measured at -20~

exhibited a maximum at a composition

of about 80 % of PC for blends of series I. Chiang and al. [85] analysed blends containing ABS of two diverse composition (AN/B/S): a) (22/19/59) and b) (20/32/48).The Brabender torque exhibited a monotonic decreased starting from the PC value with ABS addition; the PC processability improvement was larger for (b) than for (a). This difference in behaviour was attributed to a diverse morphology existing in the two systems. In other words the outer shell of the ABS particles was supposed to be made by AN in (a) and by B in (b), the latter yielding a worse compatibility with PC. This created a larger lubricating effect and hence a lower torque. Both the samples exhibited a broad minimum in the elongation at break at a different composition: 30% ABS for (b) and 50% ABS for (a). The lzod impact strength showed a pronounced maximum at a PC/ABS ratio of 90/10 for (b) and only a small one for (a). Both the effects were attributed, once again, to their different B content. Herpels and Mascia [86] used two different ABS with 20% and 30% of B in PC/ABS blends. The two kinds of blends were compared at equal rubber level in the final product. Small changes of Tg of the three phases were interpreted in terms of a partial compatibility between AN and PC. No correlation of the rubber blend content was found with fracture, obtained at low and high propagation speed. However a toughness increase was obtained under plain-strain conditions. The B was found to be encapsulated in the SAN matrix. Lombardo et al. [87], analysed over the full range of blend composition PC/ABS blends based on different types of ABS:

517 a) SAN (with no B) containing 25 % AN (SAN 25); b) an emulsion-made ABS, with 50% rubber content and very small uniform particle size (about 0.1 ~tm); c) a mass- made ABS, with 16% rubber content and large rubber particle sizes (0.5-1 ~tm). The sample a) exhibited lower modulus and tensile strength, but excellent performances in standard and sharp notched tests at R.T. as well as at low temperatures close to the Tg of the rubber. These results were attributed to the high rubber content but it was stressed that composition, rubber particle size and distribution could have also been partially responsible of the impact behaviour. The influence of the type of rubber particles was, therefore, separated from that of the rubber content. This was accomplished by making blends of SAN 25 with different ABS, in order to prepare PC/ABS blends with same rubber concentrations (5 %, 10 %, and 15 %). The results of the impact properties showed that small and uniform rubber particles (yielded by emulsion-made ABS) toughen PC/ABS materials at lower rubber concentrations and lower temperatures than large rubber particles (obtained by mass-made ABS). It was stressed, however, that also other factors, such as morphology, MW, grafting degree, agglomeration of particles, etc., can play some role in determining the above described behaviour. Wu et al. [88] studied the influence of MW of PC, in the range of M w going from 1.8x104 up to 3.6x105, on PC/ABS blend properties. They found that an increase of MW improves fracture toughness of blends at low temperature. Higher impact strengths, higher critical strain-energy release rate and lower brittle-ductile transition temperatures were, in fact, observed. However, as one would expect, the M.W. increase yields a too high melt viscosity, which renders difficult the blend processability.

518 Therefore a compromise must be reached in choosing the M.W. of PC, in order to take into account both the effects. With the highest M.W. used (3.6x105), the best compromise between impact properties and processability, relatively to the blend composition, was provided by the PC/ABS (65/35) blend.

4. Concluding remarks PC/SAN blend analysis and results can be partially utilized for PC/ABS blends, particularly with respect to component processability, miscibility and thermal behaviour. The rubber-SAN addition modifies mechanical behaviour and particularly improves, of course, the impact performances of the PC/ABS blends.

4.1 Processability A series of papers report on the lowering of the PC viscosity (or of the internal overall friction measured by the torque in mixing apparatus) by means of ABS addition, up to about 30-40% of ABS, indicating an improved PC processability [23,31,57,58,70,75,85,88]; this effect is detectable in PC/SAN blends as well [23,31] and is substantially due to a complete immiscibility of the two blend components.

4.2 Miscibility and thermal behaviour PC and ABS are completely immiscible both in melt as well as in the solid state. Several authors [23,60,63] have considered the two components partially miscible. This was made on the basis of the inward shifts of Tg detected for PC/ABS blend components (with reference to pure PC and ABS Tg values), as already done for PC/SAN systems. But for the latters this effect was clearly shown to be due to low M.W. SAN species migration towards PC boundary domains and not to partial miscibility of PC and SAN [16,27].

519 It is to be noted that these shit, s result to be larger for PC/ABS than for PC/SAN blends [23], probably because of an additional contribution to PC plasticization by low M.W. species of B. As a further evidence of this effect it was found, in PC/ABS blends containing different ABS, that the higher the B amount in the used ABS, the larger the shifts [77]. Some authors have proposed the existence of interzone layers at the PC/SAN [15,54,55] and PC/ABS [56,62,63,76] boundaries, responsible for the good adhesion between the phases; in these regions, in fact, there is the possibility of entanglement formation between PC and SAN or ABS chains.

4.3 Mechanical properties In general PC exhibits a very ductile behaviour in tensile tests on unnotched specimens. The ABS addition lowers the yield stress value, broadening more and more the relative peaks; reduces the cold-drawn ability of the PC matrix and the consequent elongation at break of the blends [59,76]. After the phase inversion, the blend curves resemble to that of pure ABS. The mechanism of deformation is mainly shear yielding in pure PC and in blends in which PC is the matrix. With increasing the ABS content the mechanism changes smoothly to crazing. The sequence of the microstructure deformation is craze initiation and propagation in the ABS, followed by arrest of crazes in the PC, and finally by extensive matrix and rubber particle deformations [ 15]. In contrast with these conclusions other authors found, by flexural [62] and by tensile tests [68], that voiding around rubber particles and cavitation, coupled with shear yielding in the PC matrix, are the main mechanisms of stress relaxation and toughening. Crazing, although precursor to final fracture, comes later on after noncrazing mechanisms and provides only a small contribution to the overall plastic deformation.

520 4.4 Impact properties Several authors have found synergistic effects in impact performances PC/ABS blends when PC is the matrix [56,61,63,66,69,76,77,82,83,85,86,88] only in same case [59] the curve of impact strength versus composition exhibited a minimum for intermediate values. From impact curves as well as by fractografic analysis [71,72,76,77] it was possible to establish that, with varying blend composition, an alternation of brittleto-ductile mechanisms occurred. Passing, in fact, from PC to blends containing more and more ABS, plain-strain to plain-stress overall transitions were observed. These effects at a microscopic level were likely due to a percolation mechanism [76] similar to the one proposed by Wu for impact behaviour of polyamides toughened by functionalized EPR rubber [89-92].

4.5 Influence of the ABS type on toughening The ABS composition is an important variable with respect to blend properties: a) The S/AN ratio is generally very close to the azeotropic S/AN ratio used for making SAN (about 25 wt % of AN) for most of ABS, whose value gives best adhesion between PC and SAN; therefore it has a relative influence; b) The PB amount is, instead of the utmost importance in determining the impact behaviour of PC/ABS blends. However literature results are somewhat contradictory: sometimes an increase of PB yielded synergistic effects [77,81]; in other cases minor improvements or even a worsening [59,82,85,86] of properties were obtained with such an increase. The reason is that a variety of parameters, other than ABS composition, can influence the impact performance: a) ABS and PC molecular characteristics [82]; b) molecular orientation; c) different processings and processing variables; d) thermal history; e) ABS rubber particle type, size and size distribution [87]; f) test

521 variables, such as temperature, strain rate, specimen thickness, notch radius, etc.; g) interfacial adhesion; h) low M.W. contents of ABS.

Acnowledgements This work was partially supported by II~ Progetto Fmalizzato Chimica Free e Secondaria of Italian CNR. The author wishes to thank D. R. Paul for sharing unpublished manuscripts on PC/SAN and on PC/ABS blends.

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527 CHA_IVrER I0

R U B B E R MODIFICATION OF BIODEGRADABLE POLYMERS

M. AveUa, B. Immirzi, M. Malinconico, E. Martuscelli, M.G. Volpe

National Research Council of Italy, Institute of Research and Technology of Plastic Materials, 80072 Arco Felice (Na) ITALY

M. Canetti, P.Sadocco, A. Seves*

Stazione Sperimentale Carta, Cellulosa, Fibre Tessili Naturali e Vegetali, 20100 Milano ITALY

1 Introduction

In order to develop materials with improved performances, copolymerization and blending have proved to be both alternative routes to introduction of new polymers.

Blending is the less expensive method.

Furthermore, while

copolymerization is confined to primary producers, polymer blending is basically a variation of compounding or formulating in general and many plastics, rubber and coating converters are well able to effect trials.

The success of such

approach is well documented by the large number of multicomponent polymeric systems introduced [ 1-3)]. As compatibility in polymers is usually an exception, different polymers, when mixed together, give rise to incompatible blends. In order to reduce the interracial tension, most successful proved to be the addition of "compatibilizers" or the induction of chemical or physical-chemical interaction during the blending

528 process [4]. This last case is quoted Reactive Polymer Blending (RPB) [5]. The main idea of RPB is to impart chemical reactivity to the polymers to be mixed, suitable to induce their compatibilization. The requirements for reactive blending in the melt consist of sufficient mixing to establish the desired morphology, presence of functionalities on both polymers and suitable reactivity across the melt phase boundary to form copolymers. Generally this kind of blending is used for polymers that contain functional groups in the backbone chains [6-7], which can undergo exchange reactions, such as amide-amide and amide-ester [8] interchange and trans-esterification under suitable reactive conditions. A different approach is when the modification is carried during the polymerization of one or both components. Interpenetrating Polymer Networks (IPN's) all fall in this last category: by a proper choice of monomers, crosslinkers and polymerization conditions, completely new products have been realized and, in some cases, starting from existing polymers. Reactive blending concurrently with the polymerization reaction is a very convenient method to obtain rubber-modified thermoplastic materials.

The

advantages of this procedure towards the melt-mixing of the components are numerous: - the polymerization of the matrix and the preparation of the blends are made in a single step, saving time and reducing the machining of the materials, which always produces some degradation; - the dispersion of the rubber component into the matrix is finer than that obtained with a melt-mixing process; -

in RPB by melt-mixing the range of the obtainable morphologies is limited,

normally, to spherical inclusion of one polymer into the other, while in the present case, unusual morphologies (core-shell, "salami" structure, etc.) are frequently encountered. Reactive blending is already applied in the case of rubber toughening of several thermoplastics.

It is found that by polymerizing caprolactam in the

529 presence of finely dispersed ethylene-co-propylene rubbers (EPR) bearing functional groups, or ethylene-co-vinylacetate rubbers (EVA), high performance polyamide 6 (PA6)-based multiphase blends with well defined morphologies are obtained [9-13].

Similar results are found for the polymerization of

dimethylterephthalate and 1,4-butanediol in the presence of dispersed reactive rubbers. More entangled morphologies are reported for the methyl methacrylate (MMA) polymerization in the presence of EVA rubbers [14] and in styrene polymerized in the presence of poly(butadiene-co-acrylonitrile) rubbers [15-16]. RPB approaches have been tempted for the toughening of naturally occurring polyesters, and the results are hereafter reported. Poly(13-hydroxybutyrate) (PHB) is a biotechnologically produced polyester that constitutes a carbon reserve in a wide variety of bacteria.

Rex:ently,

copolyesters of 3-hydroxybutyrate (HB) and 3-hydroxyvalerate (HV) have been isolated from Alcaligenes Eutrophus [17-18]:

/~H~

0II ~ #~2Hs

0II

-J~CH ,,,CI-12----C--O/~kCH-- CI-12--C ----0 HB

HV

PHBV are highly crystalline polymers with melting points and glass transition temperatures similar to polypropylene. Due to the characteristic of biodegradability (through non toxic intermediates) and processability, PHBV are being developed and commercialized as ideal candidates for the substitution of non-biodegradable polymeric materials in commodities application [19-20]. Their commercial development on such large scale is severely limited. Until recently, the prohibitive cost, narrow processability window (the difference between degradation and melting temperature), and, above all, the low impact

530 resistance around room temperature and below due to very high crystallinity and relatively high glass transition have prevented larger commercial usage. Independently, several research teams are approaching the problem of extending the potential

of PHBV through

various

blending methods.

Fundamental research has been carried out on solvent cast blending of PHB and poly(ethylene oxide) (PEO) [21]. Interesting correlation were found relating to the crystaUinity of blends at different composition. Blends of PHB and poly(methyl methacrylate) (PMMA) were obtained by melt mixing [22] and preliminary results were reported on their chemical-physical interaction. Other results on the thermodegradative behaviour of PHB-based blends [23] showed an enhancement of thermal degradation temperature by addition of poly(epichlorohydrin). When facing the toughening of PHBV, one has to keep in mind that not only the mechanical performance have to be improved, but also the biodegradability of the matrix must be saved. In this respect, RPB seems to be a realistic path: in fact, the chemical interactions are normally confined to few sites in the interfacial region, leaving unchanged the biodegradability of the matrix.

The only

condition, then, is to properly select a second polymer able to impart all the requested mechanical properties. We have used three different polymers in conjunction with PHBV: - poly(butyl acrylate), a low glass transition temperature (To rubber with a reported biocompatibility; - polycaprolactone, a low Tg semicrystalline polymer with a well recognized biodegradability [24-25]. - atactic poly(epichlorohydrin), elastomeric polymer non biodegradable. This review then deals with three different topics: 1) toughening of PHBV by "in situ" polymerization of butyl acrylate (PBA) following two different blending methods: bulk polymerization and suspension

polymerization;

531 2) toughening of PHBV by reactive melt blending with preformed polycaprolactone (PCL) in the presence of organic peroxide. 3) solution blending of PHB with atactic poly(epichlorohydrm).

System 1" toughening of PHBV by "in situ" polymerization

2.1 Bulk polymerization We have developed a method [26] in which the PHB (or PHBV) powders, as they came out from the bacterial polymerization and subsequent purification process, are thoroughly mixed with proper amount of acrylate monomers (or a mixture of acrylate monomers) and free-radical initiator resulting in a minor elastomeric phase intimately dispersed in the polyester matrix. The typical preparation of a blend is afferthere described: 70 g of polyester powder are intimately mixed with 30 g of acrylic monomer, like butyl acrylate, into which 60 nag of benzoyl peroxide (0.2 wt-% of acrylic monomer) are dissolved. The mixture is gently stirred for 24 h at room temperature. Subsequently, this homogeneous mixture is warmed to 90-100~ (extemal temperature) and left at this temperature under stirring for more 24 h, to allow acrylate polymerization. The residual monomer is extracted by vacuum stripping at 80~ and the recovered amount of the resulting blend is about 95 g. The yield of polymerization is 80-90%. It is worth to note that the inhibitor, always present in reagent grade acrylate, was not removed, in order to allow sufficient time for homogenization of reactants before polymerization. PHB homopolymer (PHBV0) and PHBV having 4 mol-% and 7 mol-% of HV comonomer (PHBV4 and PHBV7 respectively) were tested.

Although

several PHBV/PBA compositions were tested, it was found that by this method 30 wt.-% of poly(butyl acrylate) (PBA) is necessary in order to obtain relevant increase in the impact performance.

532 It was found that, occasionally, in the preparation of PHBV4/PBA blends, the acrylate polymerization occurs in regions macroscopicaUy separated from the polyester. However, at 0 %HV content, the growth of PBA occurs mainly on the PHBV powders which tend to agglomerate into larger grains. Such an effect, although not yet fully understood, could be related to change in the chemicalphysical accessibility of PHBV powder surface to BA monomer. The prepared "bulk" blends are given in Tab. 1. Sample PHBV0/PBA 70/30 PHBV4/PBA 70/30 PHBV7/PBA 70/30

% PHBV4 (g) % PBA ~g) 70 30 70 30 70 30

I I

Table 1. Composition of prepared "bulk" blends 2.2 Suspension polymerization [27-29] The bulk polymerization is scarcely reproducible, due to the lack of control of the mass polymerization of acrylic; in fact, by using an organic peroxide as polymerization initiator, like dibenzoyl peroxide, we observe an increasing of reaction temperature, that can cause the volatilization of monomers. It is mainly for the problem of exothermicity, that industrially the acrylics are polymerized in emulsion. These observation prompted us to attempt the suspension polymerization of butyl acrylate (or methyl methacrylate) in the presence of preformed PHBV powder.

By such modification we aimed to achieve a better control of

polymerization conditions.

We then prepare an aqueous emulsion of PHBV

powder to which the acrylic monomer together with peroxide is added. Under stirring, the acrylic polymerizes through a radical process induced by thermal

533 decomposition of peroxide.

The acrylic chains grow exclusively on PHBV

grains, probably due to an energetic gain. The typical preparation of this type of blend is the following: in a cylindrical reactor, 70 g of PHBV0 powder are dispersed in 400 mL of an aqueous solution containing 2.27 g of Na2HPO4, 0.13 g of NaH2PO4 and 4 g of a polyacrylic acid as emulsifying agent. The reactor is placed in a n oil bath at 110~ (extemal temperature) and 30 g of butyl acrylate, containing 0.3 g of benzoyl peroxide (1 wt-% of acrylic monomer) are added under mechanical stirring at fixed and controlled speed. The reaction proceeds for 24 h. After this time, the resulting blend is filtered, washed with boiling water and dried in vacuum. As aspected, this procedure really allows a better control of the polymerization condition in terms of heat exchange.

The acrylic polymerizes

through a radical process induced by the thermal decomposition of the peroxide. Moreover, probably, as consequence of favourable energetic balance, the acrylic phase grows exclusively onto PHBV0 chains; an important side effect of this is that the final shape of the blend particles is perfectly spherical, as normally achieved in suspension polymerization, and this is an advantage for subsequent processing of raw materials. The pol3nnerization yields are always very high (in the same range of bulk polymerization). In Table 2 we report the prepared "'suspension" blends

Sample PHBV0/PBA 80/20 PHBV0/PBA 70/30

% PHBV0 (~) 80 70

% PBA (~) 20 30

Table Z. Composition of "suspension" blends

It is worth to remark that the PHBV0 used for suspension process is different from PHBV0 of the bulk process. Particularly, they belong to different

534 bacterials straw, and this might have influences on the microstructure, and, hence, on the chemical-physical properties.

2.3 Blend characterization

In Tab 3 and 4 we report the thermal parameters of the bulk blends as obtained in differential scanning calorimetry experiments. Differential thermal analysis was carried out by using a Mettler TC 3000 differential scanning calorimeter. Two series of experiment were performed: in the first one, the samples was heated from room temperature to 200~ of 10~

cooled at the same rate down to -100~

10~

at a rate

and re-heated to 200~

In the second experiment, settled to visualize the glass transition

temperature, the sample was heated from room temperature to 200~ 20~

at

at

quenched down to -100~ and re-heated to 200~ at 20~ It is possible to observe that the transition temperatures of PHBV do not

change upon blending. Particularly Tg values are similar, thus indicating a bulk immiscibility of the two blend components. From the values reported in Tab. 3 [Part A) and Part B)] it is evident that the presence of PBA strongly depresses the enthalpy of fusion of PHBV, as well as their crystallization temperatures (particularly for PHBV0) and the enthalpy of crystallization. i lll

Sample PHBV0 PHBV4 PHBV7 PHBV0/PBA 70/30 PHBV4/PBA 70/30 PHBV7/PBA 70/30

Tm (~ I RUN

~ I * (J/g) I RUN

1740 160.9 156.0 176.3 164.8 160.3

122 97 82 86 73 57 i

Part A

T~(~ cryst

AH* (J/g) cryst

73 63 . . . . 55 60 . . . .

76 39 16 27

535

Tr (~


49.0 57.2 62.7 53.3 57.2 52.1

AH* (J/g)

Tm(~

AH* (J/g)

c~st

11 RUN

11 RUN

12 37 43 43 17 34

174.9 165.6 163.0 175.1 166.0 164.7

110 88 70 70 63 50

Part B Table 3. Melting temperature Tm,, crystallization temperature Tr and enthalpic content All of the PHBV0, PHBV4 and PHBV7 and their "bulk" blends with 30 wt-% of poly(butyl acrylate), Part A) I RUN and Crystallization, Part B) II RUN (* the values are normalized with respect to the PHBV0 content) The reduction in the ability of PHBV to crystallize upon slow cooling from the melt is a clear indication of a reduction in primary nucleation of PHBV's in the presence of PBA. This must be related also to a "reacted interface" between the matrix and the rubber, which causes a reduction in the overall crystallization rate.


T~ (~ 2.5 1.1 -0.1 2.5 1.2 1.2

Tm (~ 173 5 165 6 163 3 175 1 167.3 163 4

Table 4. Glass transition temperature Tg and melting temperature Tm of PHBV0, PHBV4 and PHBV7 and their "bulk" blends with 30 wt-% of poly(butyl acrylate).

536

2.4 Suspension polymerization Thermal analysis has been also used for the characterization of blends obtained by suspension polymerization. Surprisingly, following the same thermal treatment, rather different results are obtained. Particularly intriguing are the results obtained by the slow thermal process. It seems that the presence of 20% of PBA phase increases the enthalpy of melting and crystallization of PHBV0 phase. Similarly, in the experiments designed to obtain information on Tg's, we observe an increase in Tg of PHBV0 in the presence of PBA. As we have no reason to expect such differences from bulk to suspension polymerization process, we rather believe that the results can be at least partly attributed to the different origin of the PHBV0 used for bulk and suspension processes.

Sample PHBV0 PHBV0/PBA 80/20 PHBV0/PBA 7O/3O

Tm (~ I RUN 175.1 178.9 174.5

ii

~ (J/g) I RUN 78.0 107.1

Tr (oC) 86.1 73.4

All (J/g) c~st 64.3 75.4

Tm (~ II RUN 175.1 177.5

AH (J/g) 11 RUN 76.9 93.1

81.1

76.8

52.4

164.0

76.3

i i

i

Table 5. Melting temperature Tm, crystallization temperature Tr and enthalpic content AH of the PHB/PBA suspension blends with 20 and 30 wt-% of poly(butyl acrylate). iii

Samples

Tg PBA Tg erm (oC) , (~

PHBV0 PHBV0/PBA 80/20 PHBV0/PBA 70/30

---41.6 -43

ii

ii

Tr176 cryst

AH(J/g) cWst

Tin(~ 11 RUN

AH(J/g) 11 RUN

1.4 6.4

53 57.6

32.7 36.9

167.0 173.3

80.6 70.5

4.9

57.7

37.9

171.6

74.9

Table 6. Glass transition temperature Tg, melting temperature Tin, crystallization temperature Tr and enthalpic content AH of the PHB/PBA suspension blends with 20 and 30 wt-% of poly(butyl acrylate)

537 It is conceivable that the present PHBV0 contains either impurities or low molecular weight species which are removed upon the treatment with our reaction medium

2.5 Etching procedure Morphological characterization was carried out on the A u ~ d coated surfaces of the dumbbell specimens using Philips SEM 501 electron microscope; the samples of blends were smoothed with a LKB Ultramicrotome by means of a glass knife and etched with b ~ o l SEM

micrographs

of

smoothed

vapour in order to remove the acrylic phase. and

butanol-etched

PHBV0/PBA

PHBV4/PBA molded samples are shown in Fig. 1 and 2.

and

For the sake of

comparison, also PHBV4 was etched with n-butanol, leaving an unmodified surface. The smoothing technique was sometimes by itself sufficient to remove the rubber from its domains, especially when a not freshly made glass knife was used, as result of the shearing of the domain contour (see Fig. 2b). Of course, the use of n-butanol is necessary in order to fully reveal the morphology. It can be seen how the rubber is segregated in very small domains. No large differences were found in rubber domain size distribution between the blends containing PHBV0 and PHBV4 :.

~,

or :

....

. . . " ~ . . . . : .;~:n

,... '.....~.;~,.~=.

: , ' . : ; ' ~ , ~.

"....',,'~,;,.,..:.,.:;~

:~_ e...,

:

~e- . , - ' . , 't "

,.

r ,. ." * o~" * .?'' "

,

,',i

. " - ~', V

~'

~.": ,, ". ..:,~,~:/'... .....

~ '"'"'~.."

'~_'..,:., 9 ~ ;~ ~ ' ~ , ~

" 9 " " " # l l" " . ' . ~ t ~~ ' . ' ~ ~'" ,;~ . " ~ '9i " ""

p;..,,~.;.,.~..~. O.~,?~,~l~.~r~.l

,-',,i.

a b Fig. 1. SEM micrographs of PHBV0/PBA 70/30 blend: a) before etching, b) after etching

538

Fig. 2. SEM micrograph of PHBV4/PBA 70/30 blend: a) before etching b) after etching As it has been observed, by swelling experiments, that the permeability of PHBV to BA is very poor, it is conceivable that the polymerization occurs in the interstices between the aggregated polyester powders (the agglomeration is caused by the purification process following the biotechnological synthesis) (see Fig. 3).

Fig. 3. SEM micrograph of original PHBV4 powder

539 The post-treatments of the blends (mixing in the Rheocord chamber and compression molding) should not substantially alter the original morphology obtained in synthesis, due to the partially crosslinked nature of bulk-polymerized PBA [30].

2.6 Impact properties The impact properties were analyzed according to the Linear Elastic Fracture Mechanism (LEFM) approach [31].

The procedure used for the

calculation of the critical strain energy release rate (Gc) is reported by Coppola et al. [32]. Charpy-type specimens (6.0 mm wide and 60 mm long) were cut and notched with a fresh razor blade.

Then they were fractured at different

temperatures and at an impact speed of 1 m/s by using an instrumented Charpy pendulum. The impact properties of samples of PHBV0 and PHBV4 polymers and of their blends are reported in Fig. 4. It is evident the positive influence exerted by the rubber on the fracture toughness of polyesters. The effect is particularly marked at temperatures close and above room temperature.

In fact, while the PHBV0 and PI-IBV4

homopolymers are still rather brittle at room temperature, their blends are much more ductile. The enhancement of properties is more pronounced for PHBV0, which is known to be very brittle due to high crystallinity. The effect is less pronounced on PHBV4, probably because the addition of valerate comonomer is already effective in the induction of ductility [33]. There is a clear shill of the brittle to ductile transition in blends, compared with homopolymers. Once again, the effect is dramatic for the very brittle and stiff PHBV0.

540 Fracture t o u g h e n e s s v e r s u s T e m p e r a t u r e

,.

G (kJIm 2)

..

a

i-I PHB pure OPHBIPBA 9PHBV 4 mol % HV OPHBVIPBA

m

..

..

I

u

O-

-1 O0

I

I

-80

-60

I

- 40

T

I

I

I

l

-20

0

20

40

(oc)

Fig. 4. Fracture toughness of PHBV/PBA blends as a function of temperature By comparison of the micrographs of the blends by means of SEM, we note that the particle size distribution is quasi-bimodal, (i.e., larger spherical particles of 10 ~tm diameter co-exist with fine particles of less than 1 ~m size). It is known that, in thermoplastics technology, a bimodal distribution of a dispersed rubber is more effective in the induction of plastic deformation [34].

2.7 Graft interpolymerization Swelling experiments have been carried out on several blends in order to investigate the possible formation of cross-links between PHBV and PBA "bulk" blends.

It is, in fact, well known that the addition of peroxides to aliphatic

polyesters, like polycaprolactone, can reduce the formation of macroradicals by extraction of labile hydrogen from the polymeric backbone [5]. It is conceivable that PHBV can undergo similar reaction and so, in the presence of BA monomer

541 that PHBV can undergo similar reaction and so, in the presence of BA monomer and/or growing PBA macroradicals, intergraR processes may occur according to scheme 1.

I

I

0

0

I

H--

C--CH

I

3

I

?H 2

' C

+

R'

CH 3

I

~

?H 2

c =o

RH

+

RH

c--0

PHBV

PHBV

o

o

I

I

H--

+

C

CH 3

I

H2 C=O i

H~

+

R'

~

I

C --CH

I

' CH I C--O

PHBV

3

PHBV

macroradical (I)

I

O

I

O

-~4H9

I , C I

CH3

CH 2

I

c =o PHBV

+

o

I c--o

I

H--C --ell 2

I

C-

CH3

I

CH 2

O--C

I

!

CH 2 - - C H

I

C--O--C4H 9

II O

PHBV

macroradicals (II) The results of the swelling experiments are somewhat intriguing. In fact, by using chloroform, a common solvent for PHBV and PBA, we obtained a

542 supematant phase which accounts for about 10% of the charged "bulk" blends. DSC experiment on such residual phase did not show clear evidences of an elastomeric phase (no Tg was recorded at low temperature).

In our opinion a

possible explanation could be the following: the growth of the macroradical (II) by further addition of BA monomer is severely limited by the heterogeneity of the reacting medium. The fate of such macroradical (IT) is likely to be either the formation of short grafts or the termination on a second PHBV chain (crosslinks). On the contrary, dissolving in chloroform the "suspension" blends, we obtained much higher amount of a "graft copolymer" (about 30 wt-%), in which both transition temperatures of PHBV and PBA were found (see Fig. 5). In this case, as we found that BA chains growth only on PHBV powder, we hypothesize that this increases the possibility to obtain long graft before the PBA chain meets another PHBV chain. We hypothesize that the increased yield of grafting is due to the fact that the PBA phase grows as thin layer inside the PHBV0 powder. The improvement of phase distribution, together with the better control of polymerization kinetic, should allow the growth of longer grafted chains of PBA once they have reacted onto PHBV0 substrate.

Tg (PBA)

Tg (PHBVO)

E 0

I

-100

I

I

0

I

1

100

. . . . . . . I

I

200

Fig. 5. DSC trace of PHBV4/PBA 80/20 "suspension" copolymer

543

3 Biodegradability of PHBV To fully exploit the possible technological development of our materials it is very important to check ton which extent the substitution of 30 wt-% of PHBV with an acrylate phase influences the biodegradation of matrices. It has been reported [35] that a Gram positive bacterium which produced extracellular enzymes that degrade the homopolymer PHB when blended with the non-biodegradable atactic poly(epichlorohydrin) was isolated and tentatively assigned as Aureobacterium Saperdae. We studied the microbial degradation of plato PHBV and PHBV/PBA blends with Aureobacterium

Saperdae.

The PHBV-based blends are

characterized by mechanical and morphological methods as a function of bacterial degradation.

3.1 Enzymatic degradation Aureobacterium Saperdae cultures, where the only carbon source was the polymeric sample, were used to degrade to different extent pure PHBV4 and PHBV4/PBA blends (80/20 and 70/30 weight ratios). The micro-organism was pre-cultured overnight on 0.1% LB broth, about 3.5 mL aliquots of this culture were used to inoculate 500 mL flasks containing 100 mL of mineral medium (mineral medium composition: 1 mg/mL NH4CI, 0.5 mg/mL MgSO4.7H20 and 0.005 mg/mL CaC1.2H20 in 66 nM KH2PO 4 (pH = 6.8). The flasks were added with the polymeric samples and incubated at 30~

under shaking, for 15

days. In addition, control experiments were run to verify chemical hydrolysis of polymeric samples immersed in mineral medium at 30~

after 15 days no weight

loss of the samples was found. Cultures at different polymer degradation percentages were stopped and the samples were used to perform various morphological analysis.

544 The percentage of polymer degradation was determined by measuring the weight loss of the sample during bacterial attack.

Having checked the non-

biodegradability of PBA phase, we normalized the weight loss on the PHBV content, thus obtaining the percent of degradation in blends. Polymer samples were removed from the culture medium at different time lengths, washed several times with distilled water and dried to constant weight under vacuum. 3.2 Bacterial degradation of PHBV4 and PHBV4/PBA blends

As previously reported, Aereobacterium Saperdae is effective in the of PHB homopolymer [35]. Thus samples of plain PHBV4 and PHBV4/PBA blends (80/20 and 70/30 weight ratios) were utilised by Aureobacterium Saperdae as the only carbon source, and several polymeric specimens at different percentage of weight loss were analyzed. Preliminary tests carried out with pure PBA revealed that Aureobacterium Saperdae can not degrade this polymer. The thickness of the polymeric samples was measured before and after the bacterial attack.

0"12 !

~

~

0.1

u~ (9 C .~ r

Be

0.06

E!

0.04

=,-=

J=

P-

0.02 -

0

t

10

ii

20

t

=

I

30 40 50 Weight loss (%)

!

60

"

|

70

Fig. 6. Thickness decreasing due to microbial attack for PHBV4/PBA blends as a function of weight-loss (PHBV4 (m); PHBV4/PBA 70/30 (A); PHBV4/PBA 80/20 (D)

545 In Fig. 6 the thickness decreases are reported for the samples as a function of weight loss. As general consideration, we observed that the degradation proceeded with the same mechanism for the PHBV4 and for the blends, in fact the thickness decreases corresponded to the percentage of weight loss, indicating that the polymer erosion, by the degradative enzymes, procx~ed via surface dissolution. As further check of the exclusive surfacial erosion, test samples were cut m the transverse direction and the inner state was examined: no differences where found while degradation proceeds, for PHBV4 and PHBV4/PBA blends.

An

analogous decrease in both percentage of weight loss and percentage of film thickness, was also obtained with microbial polyesters films degraded by Aureobacterium Saperdae [35] and by Alcaligenes Faecalis depolymerase [36]. As shown by Fig. 7, plato PHBV degraded of 70% of its initial weight during 15 days of bacterial attack. 100 80 A

0

60

,o 2O

oi

I

0

3

I

I

6

9

12

15

Time (day) I-I PHBV 4 neat

II PHBV80/PBA20

~

PHBV70/PBA30

Fig. 7. Weight loss due to microbial attack for PHBV4/PBA as a function of time (PHBV4 (n); PHBV47PBA 80/20 (m); PI-IBV4/PBA 70/30 (~D)

546 Moreover, while in 80/20 blend, the bacteria consumed almost all the available PHBV4 (70 % of PHBV4 contained in the blend), the degradation stopped at 40% of available PHBV4 in the blend 70/30. The reason of such effect can be related to the morphology of blends (see below).

3.3 Morphological characterization SEM analysis of the surface of pure PHBV4 (see Fig. 8a-8d) samples after bacterial attack evidenced the homogeneous superficial erosion caused by the degradative enzymes, while no changes took place inside the sample. During bacterial degradation of the 80/20 blend, pieces of PBA component released in the culture were macroscopically visible.

As consequence of the

bacterial attack, the PHBV4 present on the surface was eroded and pieces of the dispersed PBA component were released allowing new PHBV4 zones to be accessible to the degradative enzymes

~ - ~ . . . 7

,ji;. ........

..

,,~v.~.

, ~,. ,,~ . *

.

,,,..~, .

b

i ~

:

.,2

Fig. 8. SEM micrographs of PHBV4 samples after microbial attack: a) 0% deg. (80x); b) 20% deg. (640x); c) 35% deg. (640x); d) 70% deg. (640x)

547 The SEM analysis of the degraded blend (see Fig. 9a-9d) confirms this degradation mechanism: while the degradation proceeds, large aggregates of rubber-like domains become evident, surrounded by eroded PHBV4 matrix.

Fig 9. SEM micrographs of PHBV4/PBA 80/20 sample after microbial attack: a) 0% deg. (80x); b) 23% deg (40x); c) 35% deg. (40x); d) 50% deg. (40x) The degradation, through surface erosion and PBA abiotic release, continued up to about 50% of overall weight loss. In 80/20 samples degraded at this extent the PHBV4 was no more accessible to the degradative enzymes. Anyway, such extent is very close to the amount of degradable PI-IBV4 found in the homopolymer sample test. In the case of the 70/30 blend only 20% of weight loss could be reached compared to the 50 theoretically available in 15 days bacterial exposition. No release of the PBA component was macroscopically visible in the bacterial culture and the thickness of the degraded samples did not significantly change aRer degradation.

SEM analysis (see Fig.

10a-10c)

evidenced the

impoverishment of PHBV4 zones. At the surface of the degraded blend there remains a PBA continuous phase containing small holes due to the removed PHBV4 domains (see Fig. 10c).

548

Fig.10. SEM micrographs of PHBV4/PBA 70/30 samples aRer microbial attack: a) 0% deg. (80x); b) 10% deg.(40x); c) 21% des. (40x) Such morphology do not allow a deeper penetration of bacteria, with, as consequence, an inhibition of further degradation.

3.4 Tensile behaviour

Tensile tests were performed at room temperature with an Instron machine at 10 mm/mm cross-head speed on dumbbell specimens of 1 mm thickness. The sample was cut from a compression moulded sheet prepared by heating the powder at 185~ for 5 min without pressure, then applying a pressure of 10 MPa for 2 mm at the same temperature. In Figs. 11 and 12 the trends of Young's modulus (E) and the strength at break (Ob) against the percentage of weight loss for PHBV4 and PHBV4/based blends were shown.

549

3.5 A

t~

a.

(3 = 2.5 10 O E 1.5 r

-

~

, m

.

....

1/1 t~ UJ

[3

D

0.5 i

0

....

!

I

i

I

t

I

I

10

20

30

40

50

60

70

Weight loss (%) 9PHBV 4 neat

r'IPHBV80/PBA20

A PHBV701PBA30

Fig. 11. Young's elastic modulus trends due to microbial attack for PHBV4/PBA blends as a function of weight-loss (PHBV4 (m); PHBV4/PBA 80/20 (n); PHBV4/PBA 70/30 (~) From these figures the following can be deduced: - as previously described [37] the presence of PBA particles in PHBV matrix produced a strong improvement of impact properties in the analyzed temperature range, due to the toughening effect of rubber phase (stress absorbers); however, according to above description, a decreasing of Young's modulus (see Fig. 11) and strength at break (see Fig. 12) was observed for the PHBV4/PBA blends with respect to PHBV4 homopolymers; - the modulus calculated for PHBV4 sample (Fig. 11) seemed not to vary with biodegradation process, while crb showed only a slight decreasing (see Fig. 12); this behaviour can be explained by the surfacial pathway of bacteria action, that did not allow a drastic fall of mechanical properties; -

also the reduction of Ob can be due to surfacial erosion of bacteria that can

act as crack initiator points;

550

25== a.

20-

~=

15--

A

=E L

.Q r

~p=m10~ T--5

0

0

F,

10

20

I I PHBV 4 neat

i i 30 40 Weight loss (%) I"1 PHBV801PBA20

l 50

i....... i 60 70

PHBV701PBA30

Fig. 12. Strength at break trends due to microbial attack for PI-IBV4/PBA blends as a function of weight-loss. (PHBV4 (m); PHBV4/PBA 80/20 (121); PHBV4/PBA 70/30 ( ~ ) - the PHBV4/PBA 80/20 blend, as well as PHBV4 homopolymer, presented an almost constant trend both for Young's modulus and strength at break up to 45% of weight loss; for higher weight loss (>50%), the samples seemed to loose mechanical consistency, as consequence of the increased mechanical detachment of PBA from the test sample; - finally, the PHBV4/PBA 70/30 showed a regular decreasing of mechanical properties both for the modulus and strength at break. This latter behaviour can be explained again with the morphology produced by erm3maatic degradation. In fact, the erosion of PHBV from the blend surface leads to an increase in the rubber content which for this blend is not detached from the surface, at contrary of the previous case. Consequently, the tensile mechanical properties regularly decrease with the increase in the degradation time.

551 System 2: Reactive melt blending of PHBV 4 Introduction Poly-e-caprolactone

(PCL)

and poly(]3-hydroxybutyrate-co-13-hydroxy-

valerate) (PHBV) are gaining interest from producers of polymer based goods, due to their technological properties and their inherent biocompatibility and biodegradability [3 9-44]. Since PCL can be blended with a variety of other polymers to improve their deficient properties [24-25] and PHBV is a highly crystalline polymer which is thermo-processable but brittle, a blend of the two materials could give promising results. linkage,

Since both PCL and PHBV possess the same functionality, the ester one

can,

in

principle,

employ transesterification

to

induce

compatibilization. Unfortunately, the extreme sensitivity to thermal degradation of PHBV and the high temperatures normally required for ester-ester interchange, render such a route to compatibilization impractical. In the present work we approached the problem of reactive blending of PCL and PHBV in a static mixer by means of the addition of peroxides, i.e. dibenzoylperoxide (DBPO) and dicumylperoxide (DCPO). According to the type of peroxide, two different temperatures were employed in blend preparation. For comparative purposes, the same blends had been prepared in the absence of peroxides, and the observed differences in behaviour between mechanical and reactive blends were attributed to the hypothetical formation of graR copolymer species.

Moreover the present paper is concemed with the isolation and

characterization

(thermal,

thermogravimetric

and

spectroscopic)

of the

intergraRed species from the blends prepared in presence of peroxides. Thermal and mechanical properties of all prepared blends were investigated.

552

4.1 Preparation of PCL/PHBV 70/30 with peroxide at high and low temperature Tab. 7 shows a summary of the blends and the conditions of their preparation.

The codes include the initials of the polymers, the blend

composition (weight ratio on a basis of ten) and one or two letters which indicate whether the blends are prepared at high (H) or low (L) temperature, and whether in the absence (M) or in the presence (P) of 0.5 wt.-% of peroxide. PCLPHBV 73 HP was prepared by mixing of 28 g of PCL, 12 g of PI-IBV and 0.2 g of DCPO (0.5 wt.-%) in the reaction chamber of a Haake Rheocord at a mixer roller speed of 32 r.p.m, for 10 mm at 160~

After mixing, the molten

blend was quenched by a stream of compressed air. By the same method, sample PCLPHBV 73 LP was prepared by reacting 28 g of PCL and 12 g of PHBV together with 0.2 g (0.5 wt.-%) of DPBO at 100~ For comparison, 40 g of homopolymers were mixed in the presence of 0.2 g of DCPO or DBPO in the same experinaental condition.

4.2 Preparation of PCL/PHBV 70/30 without peroxide at high and low temperature PCLPHBV 73 HM was prepared by mixing of 28 g of PCL and 12 g of PHBV in the reaction chamber of a Haake Rheocord at a mixer roller speed of 32 r.p.m, for 10 min at 160~

After mixing~ the molten blend was quenched by

a stream of compressed air. By the same method, sample PCLPHBV 73 LM was prepared by reacting 28 g of PCL and 12 g of PHBV at 100~

553 Code PHBV P PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCL P

37 37 37 37 73 73 73 73

LM LP HM HP LM LP HM HP

PCL

PHBV

T

(wt.-%)

(wt.-%)

(oc)

0 30 30 30 30 70 70 70 70 100

100 70 70 70 70 30 30 30 30 0

160 100 100 160 160 100 100 160 160 160

Peroxide DCPO DBPO DCPO DBPO DCPO DCPO

Table 7. Code and composition of the prepared blends 4.3 Extraction with chloroform 10 g of PHBV/PCL 73 P blend were treated with 500 mL of chloroform in a flask equipped with a reflux condenser and heated at 60~ under stirring for 24h. The mixture was then transferred into a separatory funnel and left to equilibrate for a few days.

The formation of two phases was observed: a bottom clear

solution and an upper phase of insoluble in chloroform.

The CHC13-insoluble

phase was recovered, washed repeatedly with CHC13, dried and then analysed. About 280 nag of solids were obtained.

The chloroform solution was

filtered, evaporated and the extract was analyzed too. By the same way 10 g of PI-IBV/PCL 37 P blend were mixed with 500 mL of chloroform to give about 100 nag of an insoluble phase.

The results of the

gravimetrical analysis are collected in Tab. 8.

Code PI-IBV/PCL 73 Res. PHBV/PCL 37 Res.

Amount of CHC13 insoluble 280 mg 100 m~;

PHBV/PCL wt. ratio in the residue 94/6 69/31

Table 8. Gravimetrical analysis of blend residuals after CHC13 extraction

554 Also PHBV P and PCL P were treated with chloroform: while PCL P completely dissolves to give a 2% (w/v) solution, PHBV P is almost insoluble, even after heating, and only swells under the action of the solvent. Pure homopolymers as well as mechanical blends dissolve completely in chloroform which is in fact a good solvent for both of them. 4.4 T h e r m a l analysis

Two series of DSC experiments were performed: in the first, the sample is slowly heated from -100~ to 200~

in the other type of experiment, designed to

study the behaviour of glass transition temperature (Tg), two heating rtms were performed at 20~

with a very fast quenching stage between them. In Tab.9

and 10, respectively, the results obtained from the two series of experiments are reported. For sake of simplicity Tab. 9 has been split in two parts" in Part A, the blends processed at 100~

are reported and compared whit the polymers, while

in Part B, the blends processed at 160~

are collected.

For comparative

purpose, PHBV, PHBVP, PCL and PCLP are included in both parts, but it must be noticed that PHBV and PCL samples are processed at 160~ and PHBVP and PCLP are treated with DCPO.

Code PHBV PI-IBV P PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCL P PCL

Part A

T.~ PCL

(oc)

37 37 73 73

LM LP LM LP

61.0 59.9 60.6 60.1 60.4 64.3

Tn~PHBV

(oc)

159.2 172.1 150.7 162.4 161.3 160.7 -

Xc,PCL

(wt.-%) 35.7 33.7 42.3 38.7 37.5 45.8

Xc,PHBV

(wt.-%) 58.0 58.8 55.0 55.7 45.6 44.3 -

555 Code PHBV PHBV P PCLPHBV 37 HM PCLPHBV 37 HP PCLPHBV73 HM PCLPHBV 73 HP PCL P PCL

Part

Tm.PCL

Tm.PHBV

(oc)

(oc)

-

159.2 172.1 161.7 171.8 164.8 160.6 -

-

60.3 61.1 60.2 60.6 60.4 64.3

Xc,PCL

Xc,PHB

(wt.-%) (wt.-%) 38.3 43.0 43.1 39.6 37.5 45.8

58.0 58.5 49.9 52.4 41.7 42.4 -

B

Table 9. Thermal and structural data of homopolymers and blends. Part A: Blends processed at 100~ Part. B Blends processed at 160~ From Tab. 9 B it is possible to observe an increase of PHBV's

Tm when

treated whit DCPO (PHBVP); such an effect is observable also in PCLPHBV 37 HP, compared to PCLPHBV 37 HM.

On the contrary the

Tm's of blends

obtained by the low temperature method are almost unaffected by the presence of DBPO. The above finding seems to suggest that some structural change occurs in PHBV when treated in the melt with DCPO. As a matter of fact, PHBV is no longer soluble in chlorinated solvents aRer DCPO treatment, indicating the formation of crosslinks. The same effect on the melting point of PHBV is found in PHBV-matrix blends, while it is absent in PCL-matrix blends, probably as consequence of the radical scavenging effect of PCL. The peroxide seems to have no influence on the PCL melting point, as well as on its solubility characteristics. In Figs. 13 and 14 crystallinity content (xo) of PCL and PHBV, respectively, is plotted as a function of the PCL percentage in the blends. A comparison of Fig. 13a and 13b shows that the blends prepared at 100~ without and with DBPO, respectively, exhibit a similar thermal behaviour, i.e. in

556 both cases x~ of PCL and PHBV decreases with decreasing their content in the blends

70

x~.(%)

X, (%)

7O

o}

60

b)

6O 5040

30

;0

'

0

20

i 60

i 80

100

30

I

I , 20

0

I

40

PCL(wt.-%)

I

60

I

100

80

PCL (wt.-%)

a b Fig. 13. Crystallmity content of PCLPHBV blends prepared at 100~ as a function of blend composition: a) without peroxide, b) with peroxide (e) PHBV, (o) PCL

Some interesting differences are found between blends prepared at 160~ without and with DCPO, as shown in Fig. 14a and 14b, respectively.

x,:(~) 70

X (%) 7O

60q

60-

4O

4O-

b)

o)

30

0

'

20

;o-

'

6o

PCL (wt.-%)

'

80

~o

1

~o0

~'---"--'------~o I

20

I

30

/

40

1

50

I

100

PCL (wt ;%)

Fig. 14. Crystallinity content of PCLPHBV blends prepared at 160~ as a function of blend composition: a) without peroxide, b) with peroxide (e) PHBV, (0) PCL

557 In fact, while for the blends prepared without peroxide the x~ values of PCL and PHBV decrease with decreasing their content in the blends (Fig. 14a), in the blends prepared with DCPO, PCL's x~ increases by decreasing its percentage in the blends (Fig. 14b). No significant differences are found between blends prepared with and without DBPO.

The origin of the above differences may be explained by a

possible compatibility of the two polymers in the molten state and the formation of grafted species between PHBV and PCL. In the above hypothesis, the blends are in homogeneous liquid state during processing at 160~

Upon cooling

PHBV crystallization may occur in mechanical blends with total rejection of PCL which is still in a liquid state. PCL then crystallizes in its own domains and is not influenced by the presence of PHBV crystals. Such a process is controlled by the crystallization rate of PHBV, which is known to be a slowly crystallizable polymer [45]. On the contrary, for the blends processed at 160~ with DCPO, the crystallization of PHBV from the homogeneous melt may not lead to total rejection of PCL phase, due to the presence of PHBV/PCL-graffed species. As a consequence, PCL crystallizes aRerwards in the interstice left by PHBV crystallization, which might well exert a positive influence on the overall crystallinity of PCL. The Tg's of the two polymers and of their blends are collected in Tab.10. Detectable variations in the Tg values are observed coveting a range of maximum 6~

At the moment, it is difficult to see a regular trend, as the

investigation is restricted to two compositions.

558 Code

%PCL

Tg.PHBV

(oc)

PHBV PHBV P PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCL P PCL

(oc)

5.1 4.1 3.2 1.8 -0.8 -2.7 0.5 3.5 3.4 - 1.0 -

-

37 37 37 37 73 73 73 73

LM LP HM HP LM LP HM HP

-62.5 -54.1 -59.4 -61.1 -54.0 -59.1 -56.4 -71.4 -58.2 -62.5

Table 10. Glass transition temperatures of polymers and blends

From the analysis of DSC traces in quenching experiments, it is possible to observe an other phenomenon: while PHBV quenched from the melt undergoes only a partial crystallization, as can be observed from the crystallization peak obtained in a second heating run (see Fig. 15a), the same polymer shows a complete crystallization in quenching upon treatment with DCPO (see Fig. 15b).

Temperature in ~

o

x ~

...... -50

.... 0

,

50

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

100 150 20 0 Temperature in ~

i

-50

.

.

.

.

i

0

.

.

.

.

i

50

,

-

-

-

I

.

.

.

.

i

.

.

.

.

i

100 150 200 Temperature in ~

Fig. 15. DSC traces of: a) PHBV (11 RUN), b) PHBV P (II RUN)

559 (B)

I -,oo

-~o

o

~o

~oo

~o

~oo

Temperolure in ~

_1oo-~'o

o

~o

a

~oo

~o

~oo

TemperQture in ~

b

Fig. 16. DSC traces of: a) PCLPHBV 37 P (II RUN), b) PCLPHBV37 P Res. (11 RUN) Blend PCLPHBV 37 HP behaves accordingly.

Once again, a possible

explanation of such results may be that DCPO produces some degree of crosslinking in the molten of PHBV. As a consequence the crystallization rate of the polyester is strongly accelerated. Thermogravimetrical experiments were performed by heating the samples from 40~ to 500~ at a heating rate of 20~

In Fig. 17 the TGA traces of

PCL and of PHBV are reported.

100

200

300

/.13.2 400 500 600 Temperoture in ~

283.8 100

200

300

Fig. 17. Thermogravimetrical traces of: a) PCL, b) PHBV

400 500 600 Temperoture in ~

560 Identical behaviour is recorded from PCL P and PHBV P. In Fig.18 the TGA pattem of the residue to chloroform extraction obtained from the PHBV/PCL 73 P blend is compared with that of the parent blend.

(A)

I/I

(:5

B

Ii

275.7 100

200

277,9

300 400 500 Temperoture in ~

100

200

300

/.00 500 600 Temperoture in ~

a

b

Fig. 18. Thermogravimetrical traces of: a)PCLPHBV 37 P res, b) PCLPHBV 37 P From the figure it seems evident that the chloroform residue is made up prevalently of PHBV with a little content of PCL, since the peak at 275~ is due to the degradation of PHBV and the shoulder at 420~

to the degradation of

PCL. TGA thus confirms DSC data about the formation and composition of copolymeric species.

4.5 FTIR spectroscopy analysis In Fig. 19 the FTIR spectra of PHBV, PCL and of the chloroform residue from the PHBV matrix blend are reported.

The spectra are recorded with a

Nicolet 5DXB spectrophotometer with a resolution of 4 cm-1. The spectrum of the residue is close to that of PHBV, bt~ the presence of a peak at 737 cm ~, due to skeletal motions of (-CH2-)5 in PCL, reveals the small, but detectable, presence of such polymer; in addition, the peak absorbance at 2941 cm 1 slightly increase

561 with respect to the one at 2988 cm 4 in PHBV, due to the contribution at 2941 cm 1 of the stretching of methylene sequences present in PCL. ,! t

/ \ t'

I--- ......

800

--I--- . . . .

750

4-- . . . . . .

700

-+---------

650

~,

600

Wovenumber

----+---'

550

I

500

.....

+

. . . . .

450

I

Z.O0

~n c m -1

Fig. 19. FTIR spectra of: a)PCL, b) PCLPHBV 37 PRes., c) PI-IBV

4.6 Mechanical properties In Tab. 11 the relevant parameters obtained from the tensile tests

of all

blends and the two polymers are reported.

Code PI-IBV PI-IBV P PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCLPHBV PCL P PCL i

E 10-3

37 LM 37 LP 37 HM 37 HP 73 LM 73 LP 73 HM 73 HP

~B

10 -2

(Ks/cm5

(Ksdcm5

20.O 16.0 14.8 12.7 15.7 11.3 4.6 5.3 4.2 5.2 2.3 3.6

2.22 3.28 2.09 1.74 2.11 1.43 1.70 1.29 1.30 1.83 4.20 5.01

Table 11. Mechanical data of the overall blends and of the pure polymers

562 The obtained data show that: -

the elastic modulus (E) decreases for PHBV- based blends with respect to

the modulus of PHBV. for the above blends, as well as for PHBV, lower values of moduli are

-

obtained when peroxides are used; E values of PCL-based blends are greater than that measured for PCL

-

which, in turn, has a higher E value compared to PCLP; -

for the above blends higher values of moduli are observed in the presence

of peroxides; -

the tensile stress at break (OB) is generally lower for blends then for

unreacted polymers, much lower in the case of PCL - based blends; -

the tendency of fibre formation in PCL is generally suppressed in its

blends. Some general explanations of the above findings are presented in the following: o the addition of ductile PCL to stiff PHBV is likely to reduce its elastic modulus.

Moreover, we have previously seen how the presence of peroxide

increases the crystallization rate of PHBV; this, likely, leads to smaller spherulites and, thus, to more ductile matrices.

The same effect may help to

explain the strong reduction of E modulus in PHBVP, compared to PHBV; - the addition of stiff PHBV to ductile PCL, on the other hand, causes a great increase in the elastic mod'~.i.i. .~--. . . ":::~

"" ....

~-.~..... " matnx " blends. The presence

of the peroxide in these blends causes a further increase in the E values. This should be attributed to the presence of grafted species at the interface between the two polymers.

We must also note that the presence of peroxide causes a

decrease in modulus in plain PCL; this might be related to degradation occurring in PCL. In view of this, the increase of moduli in the blends with peroxide seems even more relevant;

563 -

the general reduction of the tensile stress at break

((~B) in blends can be

ascribed to an insufficient interracial strength between the two polymers. This in tum, strongly depends on the amount of gaffed species produced by the radical process. Moreover, in PCL-based blends, cold drawing is suppressed and this leads, in general, to very small values of ~B. Cold drawing is still possible in PCLP, but the degradation, likely induced by the peroxide, causes a decrease in CB of the resultant fibres. The increase in ~B for peroxide-treated PHBV may be attributed to the fast crystallization of PHBVP.

4.7 Fracture analysis

The morphology of specimens broken in tensile tests has been analyzed to understand the mode and the state in which the polymeric components are dispersed and their interaction during the tensile deformation. The appearance of selected blends is shown in Figs. 20 and 21. The broken surface of PCLPHBV 37 HM and PCLPHBV 37 HP are compared in Figs. 20a and 20b, respectively.

Fig. 20. Fracture surfaces of: a) PCLPHBV 37 HM, b) PCLPHBV 37 HP

564 It is evident that, when blended mechanically, PCL disperses in spherical domains that show no interaction with the surrounding matrix, as evidenced by the smooth cavities left over by the removed PCL particles. On the contrary, evidence of plastic deformation of PCL particles is frequently encountered in peroxide-treated blends (Fig. 20b). This explains the reduction of E values found for such blends, if compared to those without peroxide previously mentioned. The appearance of PCL matrix blends is shown in Figs. 21a and 21b, respectively, where the fracture surface of tensile bars of PCLPHBV 73 HM and PCLPHBV 73 HP are pictured.

Fig. 21. Fracture surfaces of: a)PCLPHBBV 73 HM, b)PCLPHBV 73 HP

The plastic deformation of PCL, still possible to some extent in mechanical blends, obstructs the detection of PHBV particles (see Fig. 20a).

On the

contrary, the peroxide treated blend (Fig. 20b) dearly shows the PHBV particles on the broken surface, as the cold drawing of the matrix PCL is completely suppressed. The addition of the dispersed phase to the matrix polymer seems to

565 be quite high and this explains the increase of elastic modulus in peroxide-treated blends (see previous paragraph)

5 Condusions As a concluding remark, we have shown how it is possible to perform reactive blending of two bioaffine materials by the thermal decomposition of peroxides added during processing. As expected, when PHBV, a hard polymer, is present, the drawing process of PCL matrix is deterred, while the modulus of the blend increases. On the other side, PCL imparts enhanced ductility to rigid PHBV matrix. Organic peroxides

added during melt processing reduce chemical

interactions between PHBV and PCL to form a graft copolymer (scheme 2);

R' + , ' , ' ~ O ~ C

H I I

O II

~C

H2--C ~

*~ RH + , " ~ O - - C - - C

I

R'

R'

,''~O--C--

I

9

0

H

II

CH2 ~ C ~

CH3

O--C--

I

~

{~H3 O--C--

~Nww, O

+ ~

I

O

II

CH2 ~C,,*w~

C - - Cl--I2 ~C,,,,,,,,~ I II

CH 3

O

0 II Cl--I2 ~ C ' " ' ~

CH3

O II

H2-C

566

~O--C--

9

i

0 0 li i~ CI-I2 --C~w~, + ~ , , ~ 0 - - (CI-12)4 - - C H - - C ~ w ~ ,

CH~

I H

1

O

CH 3

I

~O--C--

I

II

CH2 - - C ~

~,,,, 0 - - ( C H 2 ) 4 - - C - - C ~ II H

-

-

0

Solvent extraction reveals also the presence of crosslinked PHBV; the copolymer is formed in small quantities, as the interaction in the melt is

substantially restricted to the interracial regions between the two immiscible polymers.

We can thus confirm the hypothesis that the formation of graft

copolymers is responsible for important differences found between PHBV/PCL blends obtained by simple processing and reactive processing method.

567

System 3: solution blending of PHB 6 Introduction In the present section blends between PHBV0 and atactic elastomeric poly(epichlorohydrin) (aPECH) are described, obtained by dissolution of the blend components in a common solvent.

Although no chemical interaction is

recognized for the present system, chemical-physical interactions are effective enough to give a system well entangled. Preliminary mechanical investigation, still in progress, reveals that the system is promising. The present work reports the results concerning the crystallization, thermal behaviour, optical microscopy, phase structure and influence of aPECH on biodegradation of PHBV0/aPECH blends [35, 45-46]. For thermal and phase structure investigation, the binary blends were prepared by using a poly(D(-)3 hydroxybutyrate) (Zeneca, ICI group) with a M w - 166000 and M~/M~ = 2.85, and a poly(epichlorohydrin) (Aldrich), purified by centrifugation and filtration of a

dichloromethane solution,

with

a

Mw = 1078000 and M~M~ = 3.52. The films were obtained by solution casting from dichloromethane and then dried under vacuum at 80~

Blends with weight

ratios of 80/20, 60/40, 40/60, and 20/80 PHBV0/aPECH were prepared. For the biodegradative tests, films of pure PI-IB and PHBV0/aPECH blends, which had a thickness of 40 + 5~m, were prepared by solvent-casting techniques from dichloromethane solution using glass Petri dishes as casting surfaces.

Blend

films with weight ratios of 80/20, 65/35, 60/40, 50/50, and 40/60 were obtained. The solution cast films were aged for three weeks, at room temperature, to reach equilibrium crystallinity prior to analysis [47].

7.1 Morphological and thermal characterization The thermal properties of the PFIBV0 homopolymer and blends were analyzed by using a Perkin-Elmer DSC-4 with a Perldn-Elmer 3600 Data Station (TADS System). The glass transition temperatures (Tgs) were determined by

568 heating the quenched sample from 213 to 463 K, the values were taken at the midpoint of the transition. After quenching plato PHBV0 and PHBV0/aPECH blends are completely amorphous and exhibit a single glass transition temperature (T), whose value depends on composition. The appearance of a single T, suggests the presence of a single homogeneous amorphous phase, i.e. that the two components are miscible. The experimental values show a good agreement with the theoretical Fox values [48]

1 Wpner o W,a,ecn =~ + ~ Tg T, em~ro T,~Ee,,

(1)

where Wprmv0 and T~,rmv0, and W,eECHand T~,eECa are the weight fractions and the glass transition of PI-IBV0 and aPECH, respectively. The dependence of Tg on composition is shown in Fig. 22, where the solid curve was calculated using the Fox equation (Eq. 1).

273

b.?' 263

253 I

I

I

I

0.2

0.4

0.6

0.8

1.0

Mass fraction of PECH

Fig. 22. Glass transition temperature (Tg) versus composition of PHBV0/aPECH blends. Solid curve was calculated using the Fox eq. (1)

569

7.2 Morphology and spherulite growth rate The morphology and the radial growth rate (G) values of PHBV0 spherulites were studied employing an optical polarizing microscope with an automatic hot-stage Mettler model FP 82 controlled by Mettler FP80 Control Processor.

Blend samples were melted at 458 K for 1 rain and then rapidly

cooled to the crystallization temperature (To).

During the isothermal

crystallization process, the radii of the growing spherulites were measured as a function of time by taking photomicrographs at different intervals of time. Thin films of plain PHBV0 and PHBV0/aPECH blends, when observed under the optical polarizing microscope during the isothermal crystallization processes, show birefringent spherulitic structures.

After crystallization the

samples appear to be completely filled with impinged spherulites for all the aPECH concentration studied.

The spherulite dimension, at constant To,

decreases with increasing concentration of non-crystallizable component. The spherulite radius, R, increases linearly with time for plato PHBV0 and PHBV0/aPECH blends, for all Tr investigated. For all sample the isothermal radial growth rate, G = dR/dt, was calculated at different Tr

300 A

-== 9 E E .

200

-

-

i

:L

100 -

0

358

368

378

388

398

Tc (K)

Fig. 23. Spherulite radial growth rate (g) versus crystallization temperature (To) for PHBV0/aPECH blends. PHBV0 (wt-%): (o)100; (n)80; (.)60; (0)40; (m)20

570 As shown in Fig. 23, for a given T~, the addition of aPECH to PHBV0 causes a depression of the G values, thus allowing the control of the isothermicity of the PHBV0 crystallization at lower T c values. The experimental growth rate data were analyzed acr~rdmg to the polymerdiluent theory [49-51]. The equation describing the G values of a crystallizable polymer in a one-phase melt containing a second polymer acting as a diluent assumes the following form: U* In G - In ~2 + R(T~ - r~)

_0

.2T=~ ln~b2 = a = l n G o - AK g AT r, A r f

(2)

where Go is the pre-exponential factor that includes all terms that are taken as effectively independent of temperature. energetic

contribution

of diffusional

The U*/R(T c - Too) contains the processes

of the

amorphous

and

crystallizable material to the growth rate: U* is the sum of the activation energies for the chain motion in the melt of the crystallizable and non-crystallizable molecules and Too (Too = T g - C , where C is a constant) is the temperature below which such motion cease. TOm is the equilibrium melting temperature and 62 is the volume fraction of crystallizable polymer. The term f is a correction factor that takes into account the temperature dependence of the melting enthalpy All and is given empirically by f = 2To/(Tin + To). The term Ks contains the free energy required to form a nucleus of critical size, the heat of fusion and the T~ The slope of the straight lines obtained, for plain PHBV0 and PHBV0/aPECH blends, by plotting a versus 1/T c AT f, gives the Ks values reported in Table 12. Values of U* = 10.5 kJ mo1-1 and C = 51.6 K were chosen to give the best fit least squares lines through the data.

571

_

PHBV0/aPECH

i0 "3Kg (K2)

10 7 ~~ (J an "2)

100/0 80/20 60/40 40/60 20/80

4.3 3.5 2.9 2.1 1.5

38 32 27 20 14

_

_

Table 12. Values of Ks and c~o for plain PHBV0 and PHBV0/aPECH blends In our range of crystallization temperatures, Ks can be expressed as"

K

-- nb~

(3)

AHk where n is a variable changing according to the regime of crystallization [51], b0 is the distance between two adjacent fold planes, k is the Boltzmann constant, AH is the enthalpy of fusion per unit volume, and ~ and oe are the lateral and folding surface free energies.

At the undercooling used in this study, the PHBV0

crystallizes according to regime 1II [44], then, according to the Hoffman theory (50)the n variable was set equal to 4. The values of c~e reported in Table 12, were calculated with b0 = 5.76 A [52] and c~ = 0.25 b0AH [53]. The oevalue of 38 erg -2 cm calculated for plain PHBV0 is in good agreement with the value detemainod by measurements of lamellar thickness [44]. A depression of the % value with the increasing of the fraction of the non-crystallizable component in the blend was observed in other miscible blend systems [51, 54-57].

7.3 Isothermal bulk crystallization kinetics

The weight fraction, Xt, of the material crystallized at time t was calculated using the relation:

572

[[(dH / at )at X t ._

.1o-

(4)

Io (dH / dt)dt where the first integral is the heat generated at time t and the second is the total heat when crystallization is complete.

The isotherms of crystallization of

PHBV0 and PHBV0/aPECH blends, compared at the same T~, evidenced that the overall crystallization rate constant progressively decreases by increasing the amount of aPECH in the blend. In Fig. 24 the half-time of crystallization, t0.5 (defined as the time taken for half of the crystallinity to develop), is plotted against To values for some blend composition A

w

Io o

2000

-

1500

-

1000

-

500 -

-

0

I

353

I

373

I

393

Tc ( K )

Fig. 24. Half-time of crystallization (t0.5) versus crystallization temperature (to) for PHBV0/aPECH blends (symbols as Fig. 23)

573 The overall kinetic rate constant K. was calculated by using the Avrami equation [58]:

Xt - l - e x p ( - K t" )

(5)

where n is a parameter depending on the geometry of the growing crystals and on the nucleation process. PHBV0/aPECH

T c (K)

to. 5 (s)

n

100/0

373 378 383 388 393 396 398 401 403

60 82 140 228 488 665 850 1302 2216

2.0 1.9 2.0 2.2 2.1 2.1 2.0 2.0 2.3

1.35 x 6.98 x 2.32 x 8.37 x 1.71 x 8.97 x 4.82 x 2.11 x 7.28 x

10 -4 10 -5 10 -5 10 -6 10 -6 10 -7 10 -7 10 -7 10 -8

80/20

373 378 381 383 386 388 391 393 363 368 373 376 378 381 383

319 521 605 714 882 1209 1247 2498 280 382 561 796 988 1557 1695

2.2 2.2 2.2 2.2 2.4 2.3 2.3 2.1 2.4 2.5 2.5 2.5 2.5 2.5 2.7

1.81 x 6.05 x 4.34 x 3.00 x 1.87 x 9.26 x 8.64 x 1.84 x 4.59 x 2.10 x 7.93 x 3.28 x 1.90 x 6.03 x 4.86 x

10 -6 10-7 10 -7 10-7 10 -7 10 -8 10 -8 10 -8 10 -7 10 -7 10 -8 10 -8 10 -8 10-9 10-9

348 353 358 363 368 371 373

224 360 532 673 960 1320 1636

2.6 2.9 2.6 3.0 2.8 2.8 2.7

60/40

i

40/60

K n (s -n)

2.13 x 10 -7 5.76 x 10 -8 1.95 x 10 -8 1.02 x 10 -8 3.81 x 10 -9 1.57 x 10 -9 8.69 x 10 -10

Table 13. Values of t0.5, n and K, at various To values for plato PHBV0 and PHBV0/aPECH blends.

574 =.

X !

0-

V

t... "3"

O _.1 378

1

388

393 398 I

2

403 K I

3

Log [t (s)]

I

4

Hg. 25. Log[-ln(1-xt)] versus log t according to the Avrami equation for pure PHBV0 For each To the values of n and K~, reported in Table 13, were determined from the slope and the intercept, respectively, of straight lines obtained by plotting log[-ln (1-Xt) ] versus log t (Fig. 25). The Avrami exponent, n, is non-integral with a value between 2 and 3. Similar anomalous data were also observed in the case of poly(ethylene oxide) blended with poly(ethyl methacrylate) [51].

Contrary to the theoretical

prediction [58], in almost all cases the values of n are non-integral. This fact may be accounted for by mixed growth and/or surface nucleation modes and secondary crystallization, even if this latter process does not seem to occur in our case. Experimental factors such as erroneous determination of the zero time (time when the polymer starts to crystallize from the melt), and of the enthalpy of crystallization of the polymer at a given time can cause n to be non-integral [59]

575

7.4 Melting behaviour The rate of heat evolution during the isothermal crystallization was recorded as a function of time starting on samples melted at 458 K for 1 mm and rapidly cooled to the desired Tr

After crystallization the samples were heated to the -1

melting point at a scanning rate of 10 K rain.

The observed melting

temperature (T'm) was obtained from the maximum of the first endothermic peak. The d.s.c, curves of plain PHBV0 and PHBV0/aPECH blends isothermally crystallized, showed two melting peaks, with the peak appearing at the lower temperature corresponding to the melting of the original crystal of the isothermally crystallized sample. The second endothermic peak is caused by the melting of the reorganized crystal formed during the heating process [60-61]. In fact, as shown in Fig. 26, immediately after the first melting peak, an exothermic peak is registered. The lower Tm depends on the To while the higher T~ is almost constant [60].

A . . . . . . . .

~176

I

393

I

I

413

I

I

433

Temperature (K) Fig. 26. D.s.c curves of (A) pure PHBV0 and (B) 40/60 PHBV0/aPECH blend, isothermally crystallized at 373 K

576 The

T'm values of plain PHBV0 and PHBV0/aPECH blends linearly

increase with T~ for a wide range of undercoolmg. By increasing the fraction of aPECH, a depression of the

T'm values can be observed for every To explored

(Fig. 27).

453

433

,,.-E F....

413 348

378

438

408

Tc (K)

Fig 27. Observed melting temperature (T'm) of PI-IBV0/aPECH blends as a function of crystallization temperature (Tr (symbols as Fig. 2) The experimental data can be fitted by the equation of Hoffman [62]"

m

-

yT

+

1-

m

where 7 is the morphological factor and

(6)

TOm is the equilibrium melting

temperature. In Eq. 6 1/7 assumes values between 0 ( when T'm = TOmfor all To) and 1 (when

T'm = Tr

Therefore, the crystals are most stable at 1/? = 0 and

unstable at 1/y = 1. The extrapolated

TOm value is the lower the higher the

content of aPECH in the blend is (Table 14).

577 PHBVO/aPECH

TOm(K)

100/0

461 + 3

80/2O 6O/4O 40/60 20/80

455 449 441 432

i

+2 +2 + 1 +_2

i

Table 14. Equilibrium melting temperature values for plain PHBV0 and PHBV0/aPECH blends.

The 1/3' values are very similar (about 0.3) for plain PHBV0 and PHBV0/aPECH blends, i.e. independent on composition. It is of interest to note that calorimetric measurements of other polymer crystals also yielded comparable values for 1/?. Such a result indicates that aPECH is able to act as a diluent for PHBV0 and the two polymers are miscible in the melt phase [62-63]. The melting point depression according to the Flory-Huggins theory is related to the polymer-polymer interaction parameter ~ 12 according to the relation [63-64]:

i , ll 1t

- R V 2 T;....b

T;....e + m 2 +

1t 1

m,- ~1 : f l - Z 1 2 q ~

(7)

Subscripts 1 and 2 represent the non-crystallizable and the crystallizable polymer, respectively.

All is the perfect crystal heat of fusion of the

crystallizable polymer, V is

the molar volume of the polymer unit at the

equilibrium melting temperature, rn is the degree of polymerization, To~p and T~

are the equilibrium melting temperatures of the pure crystallizable

component and of the blend, respectively, r is the volume fraction of the components in the blend and R is the universal gas constant. The following parameter values have been used in our calculation: AH = 12.6 kJ mol ~ [65]; Vl=80 cm 3 mol -1 [66]; V 2 = 76 cm 3 mol -~ [66]; m 1 =

578 1742; m2 = 7565. Using the values of TOm and T'm reported in Table 14, the plot shown in Fig. 28 is obtained. The experimental points may be interpolated by a line with an intercept at the origin (13) and a slope (X~2) of-0.068.

The

negative parameter X~2 in the PI-IBV0/aPECH system should suggest that the two components can form a compatible mixture which is thermodynamically stable above the equilibrium melting temperature.

~40 x

2-

i

2

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

(x lo)

i

.

4

.

.

.

.

--'

"-

............

I

--

6

Fig. 28. 13versus ~2 according to Eq. 7 for PHBV0/aPECH blends The fact that the line does not pass through the origin can be ascribed to the morphological effect and/or to a composition dependence of 7~2. The equation derived by Kwei and Frisch [67] for non-infinite molecular weight polymers allows the proportionality constant for the morphological contribution, C, and X~2to be calculated:

579

q~lRTm,p

k, ml J

2m2

/

- - --/~'12Zm; 1

(8)

R

Plotting the left-hand side of Eq. $ versus Ol should yield a straight line where the intercept at the origin is C/R and the slope is

-Zl2T0m,b.

The results are

shown in Fig. 29. v v

.a.

150

-

100

-

.,,...

oE

E I,..-. ! e~

E

g-

.g

50
0

ols r

1.0

Fig. 29. Application of the Kwei-Frisch equation on melting temperature depression for PHBV0/aPECH blends. The left-hand side of Eq. $ is plotted versus r

From this one determines a value of 200 for C and a value of-0.054 for Z

12"

The X12 value depends on composition and its value is similar to that derived from the Flory-Huggins equation. By comparing the two terms on the fight-hand side of the Eq. 8, one finds that C/R is greater than the term X12 T~

r

and thus

it may be concluded that morphological effects strongly influence the melting point depression of PHBV0/aPECH blends (Table 15).

580 PHBV0/aPECH

C/R

100/0 80/20 60/40 40/60 20/80

101 101 101 101 101

Z12 T~

t~1

~12 at TOm

0 -4.50 -9.22 -14.19 -19.43

-0.054 -0.055 -0.056 -0.057 -0.058

Table 15. Data from Kwei-Frisch equation (equation 8). 7.5 Structural characterization

The samples for the study of the phase structure were prepared on the basis of the thermal properties obtained from d.s.c, data.

Plato PHBV0 and

PHBV0/aPECH blends were melted at 573 K for 2 rain, then isothermally crystallized at 373 K. Then all isothermally crystallized samples were annealed for 120 rain at the annealing temperature (T) calculated as T = T' m- 19 K. The times of isothermal crystallization and the T~,w~o~are reported in Table 16. PHBV0/aPECH

t (min)

T a (K)

100/0 80/20 60/40 40/60

30 40 55 227

414 413 409 404

Table 16. Time used for isothermal crystallization at 373 K (t) and annealing temperature (Ta) for plain PHBV0 and PHBV0/aPECH blends.

The isothermally crystallized and the annealed samples, when observed under optical polarizing microscope, appear to be completely filled with impinged spherulites.

For all blends no aPECH-separated domains were

observed neither in the intraspherulitic regions nor in the interspherulitic contact zones. D.s.c. scans conducted for isothermally crystallized plato PHBV0 and PHBV0/aPECH blends, showed that for pure PHBV0 the T g of 277 K was no

581 longer detectable, while for the PHBV0/aPECH blends a single glass transition appeared at a temperature value close to the pure aPECH T g of 253 K. The T g event was more pronounced with increasing aPECH fraction.

7.6 Wide-angle X-ray diffraction The WAXD pattems were recorded on a Siemens diffractometer model D500 (Cu K~ radiation, ~, =1.54 A). The degree of crystallinity was calculated from diffracted intensity data in the range 20 = 11-40 ~ by using the area integration method [68].

The amorphous contribution was calculated from

diffracted intensity data of plain PHBV0 and blends quenched in liquid nitrogen. Lattice imperfections were not considered.

The apparent crystal sizes were

calculated from the line-broadening data collected with a scanning rate of 0.1 2 0 -1

deg min . A nickel standard sample was employed to determine the mstnunental broadening. The data reported in Table 17 show that the crystallinity x~ of the blends decreases with increasing aPECH fraction, while the crystallinity of the PHBV0 component X~,HBV0 is similar for plato PHBV0 and 80/20 and 60/40 PHBV0/aPECH blends, and lower for the 40/60 blend. i

PHBV0/ Isothermally aPECH c~stallized x~a,i-mvQ,, x~ 100/0 0.54 80/20 0.42 0.52 60/40 0.33 0.55 40/60 0.19 0.46

Annealed Xr 0.65 0.49 0.39 0.24

X~ 0.61 0.64 0.61

Isothermally crystallized (020t .(1101 208 152 271 158 289 211 284 234

Annealed (020/ 254 289 352 323

(1 i01 202 228 276 301

i

Table 17. Crystalline weight fractions x~ , crystallinity of PHBV0 component X~HBV0 and apparent crystal sizes Dhkl (A) of plain PHBV0 and PHBV0/aPECH blends.

582 The annealing produces a general enhancement of the crystallinity and the Xr

index reaches similar values for plain PHBV0 and all blends. The apparent crystal sizes Dhkl of PHBV0 in the direction perpendicular to

the (110) and (020) crystallographic planes were calculated using the Scherrer equation [68]: -

-

/80 COS0hM where 130 is the half width in radians of the reflection corrected for the instrumental broadening, and k is the wavelength of the X-ray radiation employed. The shape factor K is set equal to unity, and so the size data have to be considered as relative data [68]. The (002) reflection is disturbed by overlapping of other diffractions, and therefore the crystal dimension along the c axis cannot be calculated. The lateral crystal sizes of PHBV0, reported in Table 17, increase with increasing aPECH fraction in the blends.

This fact may be justified by the

decreased PHBV0 crystallization rate in the presence of aPECH.

In fact, the

choice of a constant temperature (373 K) enhanced the time of crystallization with increasing aPECH fraction (Fig. 14). The annealing treatment increases the above mentioned crystal dimensions and the aPECH presence seems not to interfere with the lateral increment of PHBV0 crystal dimensions.

7.7

Small-angle X-ray scattering.

SAXS data were measured at 25~ using a Huber 701 chamber [69] with a monochromator glass block. Monochromatized CuKa X-rays (X = 1.54 A) were employed. The intensity was counted at 85 angles of measurement in the range 20 = 0.1-3.8 with three different step intervals. To reduce the statistical error of counting, for each sample the mean intensity values were obtained from 16 scans

583 with a time of 21 h for a complete measurement.

The standard deviations

calculated for the intensity values at each counting angle showed a low degree of spread of the intensity data around the average values. The raw data were first corrected for sample adsorption and then the background was subtracted. By application of the indirect transformation method developed by Glatter [70-71],

the

corresponding

three-dimensional propagated

correlation

statistical

error

functions band

together

were

with

calculated

the from

unsmoothed and smeared experimental scattering data. For the desmearmg the geometries of the incident beam profile and of the detector were considered [70,72]. The scattering profiles of plain PHBV0 and some PHBV0/aPECH blends show the presence of a maximum, which is associated with the long period, L, resulting from the presence of macrolattice formed by centres of adjacent lamellae. For all S AXS measurements the abscissa variable, Q, was calculate by:

Q - 4rc

senO

2

(lo)

After desmearing the intensity were Lorentz corrected and the L values were calculated by:

27s L =

Qm where Qm is the abscissa value at the maximum of the plot (Figure 30). For the isothermally crystallized 80/20 and 60/40 PHBV0/aPECH blends, with respect to plain PHBV0 the peak position slightly shifts towards higher Q values and the peak broadens as more aPECH is added to PHBV0. 40/60 PHBV0/aPECH blend the peak does not appear at all.

For the

The annealing

produces an enhancement and a better resolution of the peak for plain PHBV0, while for blends the peak resolution does not improve at all.

584

:5 V A

0 V

0 ....I

-e-O 4-0 --e-O 4-0

!

0

,5 10 10 2 x Q ( A oI)

!

15

Fig. 30. Small angle X-ray scattering profiles of PHBV0/aPECH blends: (O) 100/0; (e) 80/20; ( , ) 60/40; (e) 40/60 In Table 18 the L values are reported for the isothermally crystallized and annealed samples. PHBV0/aPECH 100/0 80/20 60/40

Isotermall~rc~stallized 79 72 72

Annealed 93 93 82

i

Table 18. Long period distances (A) of plain PHBV0 and PHBV0/aPECH blends. A general enhancement of the long-period values can be observed as a consequence of the annealing treatment and for the blend samples the long-period distance remains slightly lower than that of plain PHBV0. The fact that the L values do not increase but are actually slightly lower for blends than for plain PHBV0 support the hypothesis that the aPECH is absent from the interlamellar PHBV0 zones. The presence of aPECH in those regions

585 could only be indicated by a very low PHBV0 lamellar thickening with blending, which is very unlikely given the crystallinity and lateral crystal dimension data obtained by WAXD.

In fact, when isothermally crystallized, the PHBV0

component reached about the same crystallinity value with blending as in the pure state.

Besides, it was observed that the lateral crystal sizes of PHBV0

increased with blending. The absence of the aPECH in the mterlamellar PHBV0 zones was confirmed by TEM measurements (not reported) which showed that the thickness of PHBV0 lamellae and of amorphous interlamellar zones are substantially the same in the plain PHBV0 and in the 60/40 PHBV0/aPECH blend. These observations together with the fact that the optical microscopy has shown the system to be completely volume filled with spherulites, and no segregation of aPECH component was observed in the interspherulitic contact zones, suggests that the non-crystallized component is segregated in the interfibrillar zones, which are larger than the interlamellar regions but smaller than the overall spherulite. To explain the scale of rejection of aPECH in the PHBV0 crystallizing matrix during crystallization from the melt, the Keith and Padden equation can be used. i.e. 8=D/G [73]. This expression places the scale of segregation on a some what quantitative basis, where 8 is the dimensional order of segregation, D is the diffusional coefficient of the non-crystallizing component in the crystallizing matrix, and G is the spherulitic growth rate. The parameter 8 has dimensions of length and represents the distance that the rejected component may move during the time of crystallization. If 8 is comparable with interlamellar distances, then the rejected component may reside between lamellae. As reported in the "Morphology and spherulite grow rate" paragraph, the presence of aPECH decreases the PHBV0 spherulite growth rate G, and the aPECH T g value is lower than that of crystallizing PHBV0. These facts give a relative high diffusion term D in the ~5parameter. The aPECH molecules diffuse away from the front of PHBV0 crystallization at such a rate so as not to remain

586 segregated between lamellae, but in any case the mobility of aPECH is not sufficient to let it move away from the spherulites. Thus, aPECH will reside in the interfibrillar zones. Considering the high aPECH M w value, a monomolecular dispersion in the interfibrillar zones of aPECH molecules with a random-coil conformation can produce scattering. In fact the volume filled by one molecule of aPECH having a weight-average molecular weight of 1078000 and a specific volume of 3

0.735 cm g

-1

has been deduced, and by the assumption of spherical shape a

radius of about 70 A was calculated. The presence of scattering arising from the dispersed aPECH molecules finds a first confirmation in the results obtained from the annealed samples. As demonstrated by WAXD investigations about the crystallinity and crystal dimensions of PHBV0, the annealing treatment produces for plain PHBV0 and blends a rearrangement and perfectioning of the crystalline region. In fact, for plain PHBV0 the profile of the Lorentz plot were better resolved after the annealing treatment (Fig. 30).

However, for the 80/20 and 60/40 blends an

increase in broadness was observed after the annealing treatment. This poorer resolution can be explained by the interference of the scattering arising from aPECH molecules dispersed in the interfibrillar zones, and this interference seems to increase with the annealing treatment. In the case of the 40/60 PHBV0/aPECH blend, for both isothermally crystallized and annealed samples the contribution to the scattering of the aPECH inhomogeneity in the system is such that the application of the Lorentz approach to the experimental scattering profile is compromised. In order to analyse the scattering arising from the presence of aPECH in the interfibrillar zones of blends the Debye-Bueche relation was used. This relation is generally applicable to the scattering intensity from the electron density fluctuation of an inhomogeneous system, it can be written as [74]:

587 [I(Q)]

-112 __

[ g 3/3 ]-1/2[1 -[- Q213]

(12)

where (n 2) is the mean-square density fluctuation in the system, 1r is the correlation length of the fluctuation and K s is a proportionality constant. For the scattering from an inhomogeneous system with correlation length 1r [I(Q)]

2

-1/2

is

linear when plotted against Q and 1'2r can be estimated from the ratio of slope to intercept. In Fig. 31 the Debye-Bueche plots for blend samples are given. It can be observed that at small Q2 values the plots are nearly linear and they deviate from linearity at larger Q2. The deviation is caused by the addition of the scattering from crystalline-amorphous region, while the lmearity at smaller Q2 represent scattering due to the inhomogeneity of the system [75].

In fact, for the

isothermally crystallized blends the lmearity continues to larger Q2 as the PHBV0 fraction decreases, giving the indication that in the blends besides the crystalline region there exist some amorphous inhomogeneity.

1

0

I

2

0

|

i!! -.5

.20 1 ",

3

I

.4

15

.3 .10

.2

.05 0

.25

.75 103 x Q2

!

1.25

0

Fig. 31. Debye-Bueche plots of PHBV0/aPECH blends. Isothermally crystallized; (e) 80/20; (~) 60/40; (m) 40/60: Annealed: (O) 40/60. The axes of this last curves are indicated by the arrows

588 In Fig. 31 the curve of the 40/60 PHBV0/aPECH annealed blend is also shown and it can be noted that the linearity continues to a Q2 value higher that for the 40/60 PHBV0/aPECH isothermally crystallized blend. The contribution of aPECH to the scattering increases as a consequence of the annealing treatment. The 1o values calculated for all blends do not significantly vary with the aPECH content or annealing treatment; the average of the obtained 1c is 45+5 A. These results indicate that no aggregation phenomena occur with increasing aPECH content in the blends. For the 40/60 PHBV0/aPECH blends, which showed a greater contribution by aPECH to the scattering a three-dimensional analysis of the scattering profile was conducted with the computation of the three-dimensional distance distribution function p(r) [70,72].

Gaussian curves were obtained, indicating

that the scattering is due to globular particles of aPECH and there was a good correlation between the scattering profiles approximated by the program and the experimental results [70]. In Fig. 32 the three-dimensional distance distribution plot for the 40/60 PHBV0/aPECH isothermally crystallized blend is reported. .8-

Q.

.4-

I

I

100

I

I

200

r(,~) Fig. 32. Three-dimensional distance distribution function p(r) for the isothermally crystallized 40/60 PHBV0/aPECH blend

589 The radius of gyration of the whole particle Rg, was calculated from the p(r) function by [70,72]:

R

~ =

*

03)

2Iop(r)d r

The Rg values obtained were about 70 A for both 40/60 isothermally crystallized and 40/60 annealed PHBV0/aPECH blends.

These values give a

measure of the dispersion of aPECH molecules with globular symmetry in the interfibrillar zones, and with the reference to the aforementioned calculated dimensions of a single aPECH molecule, aPECH appears to be dispersed at molecular level. It must be pointed out that the polydispersity of the system and the difficulties in separating the scattering that arises from the crystalline region make the experimental measure of aPECH dispersion merely an indicator of the order of magnitude of dispersion.

8 Biodegradation study PHBV0 degrading bacteria were obtained from samples of garden soil. The most rapidly growing strain was isolated by plating the cultured bacteria on agar medium containing powder of plain PHBV0 as the only carbon source.

This

Gram+ strain was characterized and tentatively designated as Aureobacterium Saperdae. Well grown A. Saperdae cultures were prepared in mineral medium containing PHBV0 as the only carbon source (mineral medium composition: 1 mg/mL of NH4C1, 0.5 mg/mL of MgSO4.7H20 and 0.005 mg/mL of CaC12.2H20 in 66 mM KH2PO 4 (pH = 6.8)). For bacterial degradation studies about 5 mL aliquots of this culture were used to inoculate 500 mL flasks containing 100 mL of mineral medium, in order to obtain 0.1 value of optical density at 540 nm (O.D.540).

To obtain comparable results, film samples

590 calculated to contain 150 nag of PHBV0 were added to each flask, then the flasks were incubated at 30~ under shaking. For the biodegradability tests, fihns of pure PHBV0 and PHBV0/aPECH blends, which had a thickness of 40+_5 ~m, were prepared by solvent-casting techniques from dichloromethane solution using glass Petri dishes as casting surfaces.

Blend films with weight ratios of 80/20, 65/35, 60/40, 50/50 and

40/60 were obtained. The solution-cast films were aged for three weeks, at room temperature to reach equilibrium crystallinity prior to analysis [47]. The

bacterial

attack

was

conducted

using

plain

PHBV0

and

PHBV0/aPECH blend films obtained from solvent-casting technique that allowed the preparation of relatively large quantities of films with homogeneous thickness.

Preliminary tests carried out with pure aPECH revealed that A.

Saperdae could not use this polymer as the sole carbon source. In Fig. 33 the O.D. 540values of the A. Saperdae cultures are plotted against the incubation time for plain PHBV0 and blends of different PHBV0/aPECH weight ratios. E

~ 1

o

~.. 1.0

v

e-

@ "0

._~

_

0

_ 0.1 Time in d

Fig. 33. Growth curves of A. Saperdae cultures where the only carbon sources was: (m) pure PHBV0; (O) 80/20; (~) 65/35; (n)60/40 and (o) 50/50 PHBV0/aPECH blends. Growth was completely inhibited with the 40/60 PHBV0/aPECH blend

591 The reported O.D.540 data are average values obtained by several culture experiments. It can be observed that the bacterial growth rate decreases in the presence of aPECH and decreases further with increasing aPECH content. Little growth occurred with the 50/50 blend and it was completely inhibited with the 40/60 blend. The decrease in the bacterial growth rate with blending could have been a result of the dispersion between the two polymers, which results in the dilution of the PHBV0 molecules on the film surface. After the stationary phase of bacterial growth was reached, the degraded films were extracted and the percentage of weight loss determined (Tab. 19). PHBV0/aPECH 100/0 80/20 65/35 60/40 50/50 40/60 i

Weig~htloss in % 100 63 37 11 3 0

Table 19. Average percentage of weight loss of blend films at the stationary phase of bacterial growth The percentage of weight loss for blend films decreased with increasing aPECH fraction, while plain PHBV0 was completely degraded.

In addition

control experiments were run to verify chemical hydrolysis of polymeric films immersed in mineral medium at 30~

After 15 d no weight loss of the films was

revealed. Since extracellular PHBV0 depolymerases have been found in bacterial broth of some microorganisms such as some bacteria (Pseudomonas Iernoignei [76], Alcaligenes faecalis TI [77], Comamonas sp. [78], and Pseudomonas pickettii [79]), and a fungus (Penicillum funiculosum [80]), cultures of A. Saperdae were assayed for PHBV0 depolymerase activity.

An aliquot of A.

Saperdae culture grown to the end of the exponential phase on mineral medium

592 containing PHBV0, was centrifuged at 10.000 rpm in order to eliminate the cells. The presence of PHBV0 depolymerase activity in the supematant was assayed at 45~

following the decrease in turbidity of a stable suspension of PHBV0

granules. The assay mixture (2.5 mL) contained supematant aliquots, 125 pmol of Tris-HC1 buffer (pH 8.0) and 2 mg of PHBV0 granules. The reaction was started by the addition of the PHBV0 granule suspension and followed by the decrease in turbidity at 650 nm, and it was calculated that 1 mL of supematant produced the degradation of 80 ~g PHBV0/mm, demonstrating that degradation still occurs in the absence of cells. In order to verify the stability of the aPECH component during the bacterial degradation, the composition of some degraded blends was determined by 1H NMR analysis. 1HNMR analyses was used to determine the composition of the biodegraded

blend samples and was carried out on a Brucker AM-500

spectrometer.

The 1H NMR spectra were recorded in CDC13 solutions of the

blends. The chemical shifts were referred to tetramethylsilane used as mtemal standard.

An average of

32 scans was accumulated, and the quantitative

analysis was carried out by the evaluation of the integrals of the signals. The results were reported in Tab. 20.

PHBV0/aPECH Weight loss PHB~0/aPECH PHBV0/aPECH initial in % b~r H NMR calculated 80/20 75 20/80 20/80 80/20 63 45/55 46/54 65/35 41 40/60 41/59 65/35 37 44/56 45/55 60/40 11 57/43 55/45 Table 20. Composition of degraded blend films at the stationary phase of bacterial growth. Assuming that aPECH was not attacked, new PHBV0/aPECH weight ratios of the degraded blends were calculated referring the % of weight loss to the initial

593 PHBV0 content of the blends. These calculated weight ratios, reported in Tab. 20, are coincident with those obtained by 1H NMR analysis confirming that during the bacterial degradation of the blends the aPECH component is untouched and/or that there is not any abiotic aPECH release. If this was not the case, the 1H NMR weight ratios would be richer in PHBV0 content than the calculated ones. Plain PHBV0 film samples degraded at different % of weight loss were prepared by stopping A. Saperdae cultures at different times, on the degraded samples the number-average molecular weight (Mn) value was determined by GPC analysis.

As shown in Tab. 21 the M n values remained relatively

unchanged during bacterial degradation.

Weight loss m% 0 22 48 61 94

Mn

Mw

Mw/M n

58400 45800 42500 45000 60000

166400 149000 142000 184000 184000

2.85 3.26 3.55 3.43 3.07

Table 21. Average molecular weights of plato PHBV0 degraded to different percentages of weight loss. This result indicates that the degradative enzymes act on the surface layer of the film and polymer erosion proceeds via surface dissolution. Analogous results were obtained by Doi et al. [81] for PHBV0 biodegradation with extracellular PHBV0 depolymerase isolated from Alcaligenes faecalis The GPC analysis of blend films revealed two partially overlapping peaks of the PHBV0 and the aPECH components. No shifts of the peaks were observed after the degradation treatment, showing that probably no important changes of the PHBV0 Mn values took place in the degraded blend samples. further evidence that no hydrolytic degradation took place.

This is a

594 The WAXD crystallinity index (x~) of solution-cast films of pure PHBV0 and blends, before and after bacterial degradation, are reported in Tab. 22, together with the crystallinity index of the PHBV0 component (X~HBV0). PHBV0/aPECH initial 100/0 80/20 65/35 60/40 50/50 40/60

Weight loss in % 100 63 37 11 3 0

Before degradation x~ x~yrmvo 0.55 0.46 0.58 0.36 0.55 0.34 0.56 0.30 0.60 0.23 0.57

Aider degradation x~ xo~,rmvo 0.54 0.28 0.60 0.26 0.57 0.31 0.56 0.29 0.60 n.d. n.d.

Table 22. Crystalline weight fraction x c and crystaUinity of the PHBV0 component x c PHBV0, for plain PHBV0 and PHBV0/aPECH blends, before and after bacterial degradation. The xr values of unattached polymer decreases with increasing the content of uncrystallizable aPECH component, while the x~, vrmv0 values are similar for plain PHBV0 and blends.

After bacterial attack, the x~ values of the 80/20,

65/35 and 60/40 PHBV0/aPECH blends are lower in direct accordance with the new PHBV0/aPECH weight ratios shown in Tab. 20 for these degraded blends They are richer in aPECH content as a consequence of the PHBV0 degradation. In contrast, the xr vnav0 values did not change in the degraded blends. Therefore, it may be concluded from the WAXD data that the weight loss of the blend films is solely due to the loss of the PHBV0 fraction in agreement with the assumption made above. The SEM analysis (not reported) of the film surfaces revealed surface erosion of plain PHBV0 and PHBV0/aPECH blends after bacterial attack. The film thickness was determined for plain PHBV0 and 80/20 blend before attack and after a 60% of weight loss. A corresponding decrease of plain PHBV0 film thickness of approximately 60% was measured indicating a surface erosion caused by the degradative enzymes. The observed decrease in both percentage of

595 weight loss and percentage of film thickness, was also obtained by Doi et al. [81] with PHBV0 films degraded by Alcaligenes faecalis depolymerase. A different result was observed for the 80/20 PHBV0/aPECH blend: aider the attack the film thickness remained unchanged, while large cavities (several ~tm) were observed inside the film. Highly degraded 80/20 film (where only 25% of the PHBV0 initially present in the blend is left) maintains the initial dimensions probably due to the aPECH dispersion in the blend. In fact, the composition analysis of the degraded blends revealed that there were neither degradation nor abiotic release of aPECH during the bacterial attack. The intemal large cavities observed in the degraded blend probably could be zones where PHBV0 was extensively degraded and where aPECH was collapsed onto the remaining structure after bacterial attack.

9 Conclusions

Miscibility in the melt between PHBV0 and aPECH is confirmed by the detection of a single glass transition and by the influence of aPECH on the spherulite growth rate and on the overall crystallization rate. The depression of the equilibrium melting temperature appeared to be strongly influenced by morphological effects. After PHBV0 crystallization from the melt the aPECH molecules are rejected into the interfibrillar zones where they probably assume a random-coil conformation. In spite of the molecular dispersion of aPECI-I, the detection for the crystallized blends of a glass transition temperature value close to the one of aPECH, indicates weak interactions at the segmental level between the two polymers.

The trend of aPECH molecules to assume a random-coil

conformation together with the high molecular weight can hinder the close contact between aPECH and PHBV0 molecules in the continuous interfibrillar zone.

The annealing treatment promotes a general perfec~ionmg and

rearrangement of the sample morphology, enhancing the crystallinity and the

596 crystal dimensions of PHBV0 in the pure state and in the blends, and probably favouring the trend of the aPECH molecules to assume a globular conformation. A. Saperdae is a Gram positive bacterium which extracellulary degrades PHBV0 blended with aPECH.

The biodegradability of blended PHBV0

decreased but was not completely compromised by the presence of aPECH component. Bacterial attack of blends containing relatively low aPECH percentages causes a drastic reduction of the PHBV0 content, while the dimension of the film remained unchanged due to the aPECH dispersion inside the blend. In fact, as a consequence of bacterial attack neither degradation nor abiotic release of aPECH was observed.

References

1. Sperling, "Interpenetrated Polymer Networks and Related Materials", Plenum Press, New York (1981) 2. Yu.S. Lipatov and L.M. Sergeeva, "Interpenetrated Polymer Networks", Naukova Dumka, Kiev (1979) 3. D. Klempner, Angew. Chem, 90, 104 (1978) 4. D. Klempner and D.C. Frisch, eds., "Polymer Alloys II", Plenum, New York (1980) 5. M. Narkis, "Reactive polymer blending by free radical processes", in "European Regional Meeting", PPS, September 1991, Palermo (Italy) 6. F. Pilati,, E. Marianucci, C. Berti, J. App. Polym. Sci., 30, 1267 (1985) 7. L.H. Wang, Z. Huang, T. Hong, R.S. Porter, J. Macromol. Sci., Phys., B29, 155(1990) 8. L.Z. Pillon, L.A. Utracki, in "Ecole d'Application des Hautes Polymeres, 3eme Conference sur la reactivite chimique des polymeres en course de transformation", Strasbourg, France (1984), p. 57

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599 44.P.J. Barham, A. Keller, E.L. Otun, P.A. Holmes, J. Mater. Sci., 19, 2781 (1984); P.J. Barham, ibid, 19, 3824 (1984) 45.E. Dubini Paglia, P.L. Beltrame, M. Canetti, A. Seves, B. Marcandalli and E. Martuscelli, Polymer, 34, 996, (1993) 46.P. Sadocco, M. Canetti, A. Seves and E. Martuscelli, Polymer, 34, 3368 (1993) 47.S. Bloembergen, D. A. Holden, G. K. Hammer, T. L. Blum and R. H. Marchessault, Macromolecules, 19, 2865 (1986) 48.T.G. Fox, Bull. Am. Phys. Soc., 2, 123 (1956) 49.J. Boon and J. M. Azcue, J. Polym. Sci. A2, ,6,885 (1968) 50.D. Hoffman, Polymer, 24, 3, (1982) 51.S. Cimmino, E. Martuscelli, C. Silvestre, M. Canetti, C. de Lalla and A. Seves, J. Polym. Sci., Polym. Phys. Edn, 27, 1781, (1989) 52.M. Yokouchi, Y. Chatani, H. Tadokoro, K. Teranishi and H. Tani, Polymer, 14, 267 (1973) 53.J.D. Hoffman, Polymer, 26, 803 (1985) 54.E. Martuscelli and G.B. Demma, "Polymeric Blends: Processing Morphology and Properties", Plenum Press, New York, 1980 55.E. Martuscelli, M. Pracella, Wang Ping Yue, Polymer, 25, 1097 (1984) 56.S. Cimmino, E. Di Pace, E. Martuscelli, C. Silvestre, Makromol. Chem Rapid Commun., 9, 261 (1985) 57.E. Martuscelli, C. Silvestre, C. Gismondi, Makromol. Chem., 186, 2161, (1985) 58.J. Avrami, J. Chem. Phys., _7, 1103, (1939) 59.D. Grenier and E. R. Prud'homme, J. Polym. Sci., Polym. Phys. Edn, 18, 1655(1980) 60.H. Mitano, P.J., 19, 1241, (1987) 61.T. Arakawa, F. Nagatoshi and N. Arai, J. Polym. Sci., Polym. Phys. Edn., 7, 1461, (1969)

600 62.D. Hoffman and J.J. Weeks, J. Res. Natl. Bur. Std., 66, 13, (1962) 63.T. Nishi, T. Wang, Macromolecules, _8, 909, (1975) 64.P.J. Flory "Principles of Polymer Chemistry", Comell University Press, Ithaca 1953 65.P. Greco, E. Martuscelli, Polymer, 30, 1475 (1989) 66.D.W. Van Krevelen, "Properties of Polymer", Elsevier, Amsterdam, 1990 67.T.K. Kwei and H. L. Frisch, Macromolecules,, 11, 1267, (1978) 68.L.E. Alexander, "X-ray Diffraction Method in Polymer Science", WileyInterscience, New York, 1969, p. 137 69.E.Schnabel, R. Hosemann and B. Rode, J. Appl. Phys., 43, 3237 (1972) 70.O.Glatter and O. Kratky, "Small Angle X-ray Scattering", Academic Press, London, 1982, p. 119 71. O. GlaRer, J. Appl. Cryst., 13,577 (1980) 72. O. GlaRer, J. Appl. Cryst., 10, 415 (1977) 73.H.D. Keith and F. J. Padden, J. Appl. Phys., 35(4),1270, (1964) 74.P. Debye and A. M. Bueche, J. Appl. Phys., 20, 518, (1949) 75.S. Nojima, Y. Terashima and T. Ashida, Polymer, 27, 1007, (1986) 76.K. Nakayama, T. Saito, T. Fukui, Y. Shirakura, K. Tomita, Biochim. Biophys. Acta, 827, 63, (1985) 77.T. Tanio, T. Fukui, Y. Shirakura, T. Saito, K. Tomita, T. Kaiho, Eur. J. Biochem., 124, 71, (1982) 78.D. Jendrossek, I. Kajoke, R.B. Habibian, A. Steinbuchel, H.G. Schlegel., J. Environ. Polym. Degrad., 1, 53 (1993) 79.K. Yamada, K. Mukai, Y. Doi, Int. J. Biol. Macromol., 15, 215, (1993) 80.C.L. Brucato, S.S. Wong, Arch. Biochem. Biophys., 290, 497 (1991) 81.Y. Kumagai, Y. Kanesawa and Y. Doi, Makromol. Chem., 193, 53 (1992)

601

Subject index

Accelerator 27

Anhydride conversion 29

Accommodation coefficient 177, 178

Arrhenius factor 30

Acetic anhydride 76

Atactic poly (epichloridrin) 532, 534,

Acrylate monomers 535

547

Acrylonitrile-Butadiene-Styrene

Aureobacterium Saperdae 547, 593

copolymer (ABS) 469, 473

Average particle diameter 304

- composition 509, 510, 512, 513,

Avrami equation 165, 166, 577

514, 515, 516, 517, 520 -

-

emulsion-made 517

Benzoyl peroxide 97, 551

mass-made 517

Biocompatibility 532, 555

- type 520

Biodegradability 530, 532, 547, 555,

Activation energy 30, 300, 302, 306

595, 601

Adhesion 54, 57, 69, 94, 290, 293,

Bismaleimide resins 123, 124, 135

294, 296, 317, 330, 335, 371, 375,

Blend 125, 184, 190, 192, 194, 197,

391,359, 389, 430, 433,476, 478

201, 202, 439, 470, 471, 492, 518,

Alloys 470

538

Amide

- binary 352, 359, 374, 380, 406,

-

- bands 64, 78, 79 -

418, 433

absorption 77

linkage 76

- components 471 -

81, 97, 106, 291, 292, 293, 296,

Amino -

-

323, 329, 495, 496, 497, 498,

groups 75

503, 507, 556, 560, 576

terminated butadiene-acrylonitrile copolymer 75, 76

composition 16, 46, 50, 52, 68,

Blending 243, 244, 254, 291 -

conditions 290, 321,330

602 - mechanical 473

Composites 479, 484

-

reactive 470, 528, 534, 554

Compressive yield stress 88, 107

solution 534, 571

Conversion of alfllydride groups 105

time 246, 263

Copolymer

-

-

Bragg equation 275

- graft 335, 349, 353, 369, 387, 399, 413, 429, 433, 464, 465,

Brittleness 124

466 Calibration

random 350, 432

-

- curve 98

Copolymerization 124, 527

- factor 35, 38

Core

Carbonyl -

-

-

shell particles 57

-

absorption 64

-

stretching 20

Crack, 455

terminated butadiene-acrylonitrile

- arrest line 41

copol3aner (CTBN) 75

- growth 34

Carboxylic groups 293 Cavitation 14, 72, 74, 93, 261,262, 317, 499

structure 439, 450, 466

initiation 507, 508, 513, 519

-

pinning 56

-

propagation 519

-

Chain scission 25, 27

- tip blunting 44

Chemical

Craze 294, 295, 317, 327, 328, 447,

fonnulas of epoxies 11

-

-

interactions 89, 143

Co-continuity 484 Cold drawing 312, 315, 316, 317, 318 Compatibilizing agent 469 Compatible 414, 429

455 initiation 480

-

propagation 519

-

-

tip 480

Crazing 498, 519 -

mechanism of 519

Critical

603

-

-

strain energy release rate (Gr 36,

Defonnation mechanism 72, 73, 488

52, 68, 69, 90, 92, 109, 125, 152

Diblock copolymer 106

stress intensity factor (Kr 35, 51,

Dicumil peroxide 551

68, 69, 80, 91, 108, 125, 151

Dielectric properties 497

Cross-Bueche equation 249, 255,

Differential

264, 269, 297, 298, 306

(DSC) 16, 46, 462, 463, 464, 495,

Cross linking

534

-

calorimetry

Dispersion coarseness 251,292, 302,

agent 440

- density 13,442, 443

317, 328, 329

Crystalline

Distribution

- forms 307

- anysotropic 292, 301,309

- lamella thickness 273, 275, 283

- layered 310, 314, 316

Crystallinity index 273, 276, 559,

-

Crystallization 157, 433

-

- conditions 243, 244, 245, 272,

-

-

of the particles 54, 57

Domain

599

size 442, 443 structure 443

282, 290, 307, 330

Droplet 171, 172, 192, 215

fractionate 214, 215

Dugdale model 106

- process 262, 272, 273, 282, 283 -

scanning

rate 281

Dynamic mechanical -

analysis 32, 50

Curing process 16, 27, 48

- behaviour 483,479, 497

Debye-Bueche relation 586, 587

Elastic modulus 126, 131, 152, 312

Decarboxylation 25

Elastomer 335

DDS

(4,4'

diamino

diphenyl

- at break 317, 318

sulphone) 18 DDM

(4,4'

methane) 12

Elongation

diamino

diphenyl

- at fracture 485 Emulsifying agent 74, 106, 109

604 Engineering thermoplastics 14

Filler 178

Enzymatic degradation 543, 550

Films 185, 186

Epoxy

Flame retardants 498

-

-

equivalent weight 13

Flexural strength 133

resins 11, 12, 13

Flory-Huggins theory 577

Etching 127, 500, 501,502, 503

Flow mechanism 298

Ethylene-propylene copolymers 205,

Fourier

206

Spectroscopy (FTIR) 16, 18, 64, 76,

Ethylene-Propylene

Transfonn

Infrared

136, 560

monomers

(EPM) 336, 340, 346, 365, 380,

Fox equation 568

399, 414

Fractographic

-

g - SA 348, 357, 369, 378, 397,

328, 454

417, 425, 337, 344

Fractography 499, 500, 507

- g - SA- PA6 349, 369, 380, 391, 429 Extrudate swell 482, 494 Extruded sheets 496 Extruder 194, 195

analysis

260,

294,

Fracture 479, 498, 499, 506, 507 -

analysis 32, 563 behaviour 34, 51, 118, 125

-

-

energy 133 fast 325

-

-

double screw 305

-

front 295

-

single screw 321

-

induction 260, 261,294, 295, 326

Extrusion -

capillary 321

-

chamber 321

-

direction 324

mechanics tests 100

-

-

mechanism 262, 296, 327, 352, 390

-

surface 72, 260, 261, 294, 325, 326, 430

Failure -

mode of 480

Fibre rupture 316

-

tougheness 12, 13, 36, 68, 69, 131

605 Glass transition temperature 12, 14,

- behaviour 245, 254, 263, 271,

18, 32, 47, 50, 246, 253, 350, 351,

328, 330, 366, 395, 398, 400,

478, 482, 490, 491, 495, 512, 513,

414, 416, 430, 496, 504, 507

529, 535, 554, 567, 595

- modifiers 262

Grafting

- perfomlance 496, 513, 520

- degree 97, 109, 341, 366, 409,

- properties 253, 262, 267, 269,

417, 424, 434, 435, 517 - of maleic anhydride 338, 420,

-

271,294, 321,539, 549 - strength 253, 259, 261,262, 263,

428, 429, 432

268, 283, 294, 295, 296, 325,

of unsaturated molecules 337,

326, 329, 470, 489, 498, 506,

345, 400

514, 520 -

test 38, 92, 125

Hardener 12, 17

- tougheness 38, 43, 51, 52

Heat

Incompatible blends 426, 438

-

-

distortion temperature 470

Injection moulded 246, 247, 258,

of cure 48, 49

266, 270, 272, 276, 277, 291, 292,

High impact polystyrene (HIPS) 443,

301, 303, 305, 307, 311, 313, 314,

452

315,316

Hoffinan theory 571

Interaction

Hydrogen bond 21, 63, 65, 66, 77,

- energy 476, 487

78 -

Hydroxyl terminated -

parameter 476, 477, 482, 496

Inter

polybutadiene rubber 61 -

- PC24 -

fibrillar regions 276 lamellar amorphous layer 276, 277,278

Image analysis 113 -

Impact

particle distance 254, 262

- penetrating

polymer

(IPN) 440, 449

network

606 -

spherulitic

amorphous

Long spacing

contact

2

7

5

Low profile additives 61

regions 279 Interface 69, 74 Interfacial

Maleic anhydride 76

-

adhesion 74

Maleimide groups 85

agents 469

Maleimido

-

-

-

fracture tougheness 490

-

tension 487, 488, 489

-

-

-

-

Maximum stress at fracture 485

terminated polybutadiene rubber

Mechanical -

-

Kinetic

-

-

terminated rubber 75, 136

groups 62, 68

62

-

temainated butadiene-acrylonitrile copolymer (ITBN) 75

Irwin model 95, 110 Isocyanate

double bonds 136

analysis 30, 32, 136

behaviour 469, 480, 504, 519

-

degradation 336

-

properties 349, 366, 393, 419, 530, 549

constant 31 equation 30

analysis 32

-

resistance 470, 505

Kwe and Frisch equation 578

Melt

Kyotami equation 307

- mixing 243, 291, 321, 335, 348, 367, 375, 398, 418, 432

Linear Elastic Fracture Mechanics

-

rheology 245, 269, 290, 321

(LEFM) 125, 458

-

viscosity 472, 492, 493, 499

Liquid

Melting

-

crystalline 218, 219

-

rubber 13, 52, 62, 75

Logaritlun rule of mixtures 247

-

point 157, 174

- temperature 529, 535, 570, 575, 595

607 Methyl nadic anhydride 12, 27

- mass 244, 245, 249, 250, 253,

Michael reaction 123, 124

254, 255, 258, 262, 263, 264,

Micro

266, 267, 268, 269, 270, 277,

- cracking 56

278, 282, 283, 285, 289, 290,

-

292, 296, 298, 300, 301, 305,

defomlation 481

307, 318, 321,329, 330, 332

- voids 371 Migration 211, 212, 213, 214, 483,

-

mass distribution 244, 245, 254,

491,513, 518

255, 263, 264, 267, 268, 269,

Miscibility

270, 282, 283

- in PC/SAN blends 478, 481,487, 518 -

window 477

- structure 245, 261,262, 272 -

weight 62, 75

Morphological analysis 40, 110, 127,

Mixing

452, 537, 546

- teclmiques 349, 350, 353, 429

Morphology 70, 349, 352, 361, 371,

- torque 503, 510, 511,512, 516

386, 393, 407, 416, 426, 432, 483,

Mode and state of dispersion 244,

484, 496, 497, 498, 502

245, 247, 258, 262, 263, 266, 267,

-

droplet-like 259

270, 271, 282, 292, 301, 315, 323,

- melt 243,263, 282

329, 330

Multicraze formation 262

Modulus 14, 36, 37, 53, 85, 106 -

-

loss 247

Necking 316

storage 247

Negative deviation blends 247, 249,

Mold 291,292, 305

270

Molecular

Notch sensitivity 470

- characteristics 290, 297, 305

Nuclear Magnetic Resonance (NMR)

-

-

composite 218 interactions 21

492, 592 Nucleating agents 176, 177, 197, 212,213

608 Nucleation -

dual continuity 442, 443

-

density 185, 186, 189, 192, 197,

inversion 55, 442, 443, 445, 449,

-

200, 201,204, 244, 283 -

heterogeneous

450, 452, 453, 454, 503

161, 163, 164,

173, 174, 189, 193, 197, 203,

- morphology 323 -

205, 213, 220 - homogeneous 161, 163, 164, 189

e

p

a

r

a

t

i

o

n

50, 376, 379, 408,

421,440, 443 -structure 244, 262, 269, 277,

- primary 158, 163, 164, 168, 170,

282, 287, 289, 290, 318, 321,

173, 184, 185, 186, 188, 192, 195, 202, 204, 205, 207, 217

s

325, 329, 330 -

viscosity ratio 252, 258, 263,

- secondary 179, 180, 182

267, 270, 271, 282, 283, 292,

- self 175, 177, 197, 203

296, 302, 303, 304, 312, 329, 330

Optical -

-

-

bistability 446 density 589 microscopy 480, 481

Photostability 496 Plane -

strain 499, 506, 509, 520

- stress 499, 506, 509, 520

- properties 446

Plastic

- transparency 439, 450

- deformation 56, 295, 325, 327

Orientation 214, 261 Oscillatory shearing flow 246, 254

-

-

shear deformation 72, 94, 111 zone 42, 95, 96

Plasticizer 48, 50 Particle size 72, 251,254, 262, 263,

Poisson ratio 91

268, 270, 271, 282, 283, 292, 301,

Poly

311,328, 329

- acrylonitrile (PAN) 470, 472

-

distribution 267

Phase

- amide 336, 348, 351, 365, 376, 400, 417, 432

609 -

-

b

i

s

p

h

e

n

o

l

-

A

carbonate 15, 16,

-

469, 471

-

butadiene 99, 444, 470

-

- butylacrylate 530, 536 -

-

-

vinylacetate 62, 464, 465, 466

butylacrylate-styrene 439, 448

Polymerization 335, 375, 400, 409,

caprolactam 419, 421

413, 418, 432, 474, 474, 474, 527,

condensation 375, 378, 418

530, 530, 536 -

shrilukage 61

Processability 470, 472, 504, 516,

- ethersulfone 15

517,518

- ethylene, high density 336

-

sulfone 15

Polymer-diluent theory 570

- ether imide 15, 46, 124

-

urethane 441,447, 448

Processing 244, 287, 290, 321, 330,

ethylene-vinylacetate 439

483, 486, 492, 497, 499, 483, 486,

esterresins 61, 75

492 -

13-hydroxybutyrate-co- conditions 321,327, 329, 330

hydroxyvalerate (PHBV) 529 -

-

-

- cycle 291,305

imides 121,122

Pseudoplastic

isobutylene 96

- flow 250

isoprene 99

- methyl methacrylate 61,439 - propylene 336, 341 -

-

-

- melts 247, 264, 297

Radial growth rate 272

styrene 470

Radical polymerization 123

styrene-butadiene 444

Reaction

styrene butadiene-styrene 445 -

-

styrene-co-acrylonitrile

(SAN) -

470, 471, 472, 473, 474, 476, 481, 484, 486, 488, 489, 490, 491,473, 474, 472, 477, 520 -

styrene-methacrylic acid 442

mechanism 101 time 69

Reactive elastomers 124, 131 Refractive index 439, 449, 466 Relaxation time 298, 302 Reptation 162, 180

610 Residence time 321, 322, 323, 326,

Spherulite

327, 328

-

average dimensions 278

Rheological behaviour 485, 510, 511

-

Rubber

- growth rate 166, 569, 585, 595

-

-

particles 71,487, 488, 517

-

in blends 389, 397, 417

phase volume (RPV) 444, 446,

-

radius 188, 199, 219, 222

447

size 165, 190, 193, 201

-

SAXS profiles 273, 275, 276, 584 Scamling electron microscopy (SEM) 41, 53, 70, 83,452, 457, 500 -

analysis 250, 258, 266, 276, 292, 309, 312, 323

Selective dissolution 324 Shear bands 480

-

lips 499

-

-

-

dimensional 470, 472, 473

-

thennal 471

Stick-slip propagation 34 Straining 315 Stress -

at break 317

-

concentration 317

-

cracking 470, 473 strain curves 87, 107, 505

-

-

thi~ming 249, 250, 257, 298

-

yielding 14, 72, 74, 111, 262, 294, 295, 327, 519

Sherrer equation 582

size distribution 244

Stability

rate 247, 249, 296, 298, 321,493 stress 301,323

-

fibrillae 275

-

whitening 260, 261, 262, 295, 317, 325, 328

Structure -

interphase 272

-

layered 258, 266, 267, 309, 312,

Shish kabab 176

313,315

Size of the particles 54, 57 -

rod-like 325

-

super-reticular 272

Slow-growth region 41 Specimen 472, 481, 481 Spectral subtraction 20, 138, 144

Succinic anhydride groups 97

611 Surface - energy 209 -

free energy of folding 274, 278

System -

-

Thermal properties of PC/ABS blends 504, 512, 518

-

of PC/SAN blends 480, 482, 518

-

Thermoplastics 124

incompatible 469

Thermoplastic modifier 15

multiphase 470

Tie molecules 244, 280, 281 Toluene diisocyanate 62

Tan 5 32

Toughening 245,254, 259, 262, 263,

Taylor-Tomotika theory 252, 254,

280, 282, 283, 290, 296, 318, 329,

259, 263, 271,303, 329

330, 375, 401, 435, 439, 440, 443,

Temperature

447, 448, 449, 465,469, 470, 500

- apparent melting 272 - equilibrium melting 272

-

-

agent 13, 136 mechanism 56, 497

- extrusion 302, 322, 323, 324

- polyamide 365,377, 417

- melting 290, 291,297, 305, 332

Transparency 470

- mold 246, 263

Transmission electron microscopy (TEM) 80, 444

- processing 324, 325 Tensile -

Triblock copolymer 66, 69, 72, 74

elastic behaviour 312, 584 Ultimate tensile strength 462

- properties 317, 351, 352, 363, 370, 380, 416, 433 -

test 479

Ternary blends 486, 488 TGAP (triglycidyl epoxide based on

Undercooling 243, 272, 278, 281, 282, 283,325,329 Uniaxial compression 86 Urethane - carbonyls 65

aminophenol) 11 -

TGDDM (tetraglycidil epoxide based

linkages 63

on diamino diphenyl methane) 11, 46 Vicat temperature 496

612 Vinylacetate

content

of

EVA

copolymer 289, 290, 293, 305, 306, 307, 308, 311, 312, 315, 316, 317, 318, 320, 327, 330 Viscosity 247, 249, 251, 252, 254, 255,257, 258, 482, 494, 499 - apparent 249, 297, 322, 323 - complex 246, 247, 264, 271 - dynamic 246 -

-

Mooney 246 ratio 483 zero-shear 249, 269, 298

Viscoelastic properties 484 Volume of flow element 301,306

WAXS patterns 307, 581 Welding 500

Yield 352, 371 -

-

point 314 stress 14,490

Yielding 260 -

behaviour 86

- mechanism 43, 56 - process 326 Young elastic modulus 426

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