matrix transcrystalline interphase in carbon fibre reinforced j-polymer microcomposites

Composites Science and Technology 47 (1993) 43-50

THE MECHANICAL ROLE OF THE FIBRE/MATRIX TRANSCRYSTALLINE INTERPHASE IN CARBON FIBRE REINFORCED J-POLYMER MICROCOMPOSITES Silva Incardona, Claudio Migliaresi Department of Materials Engineering, The University of Trento, Mesiano 38050, Trento, Italy

H. Daniel Wagner Department of Materials and Interfaces, The Weizmann Institute of Science, Rehovot 76100, Israel

Adrian H. Gilbert & Gad Marom Casali Institute of Applied Chemistry, Graduate School of Applied Science and Technology, The Hebrew University of Jerusalem, 91904 Jerusalem, Israel (Received 14 October 1991; revised version received 7 April 1992, accepted 29 May 1992) third, relatively thick, intermediary phase present between the constituents. Its elastic and mechanical properties are specifically designed to produce a certain effect on the overall performance of the composite material. For example, a gradient modulus interphase approach has been proposed to promote the mechanical properties and the fatigue life by reducing the modulus mismatch of the fibres and the matrix, 1'2 while a soft, elastic interphase has been advocated for high fracture toughness, and for reduced stress concentrations in the matrix around the fibres. 3'4 A number of theoretical models link the properties of the interphase and its relative thickness to the performance of the composite material. This can be exemplified by studies of the mesophase, a region with variable properties between the fibre and the matrix) The advantages of intentional addition of an interphase, with its wide spectrum of design parameters, emerge clearly compared with a fibre/matrix interface. The latter, which is formed spontaneously when a composite material is prepared, has a single design parameter, namely, the fibre/matrix bond strength. In semicrystalline thermoplastic polymer matrices, crystallization in the presence of fibres may result in the development on the fibre surface of a transcrystalline region which acts as an interphase. This has been shown for a number of polymers, including PEEK, 6 J-Polymer,7 polyphenylene sulphide 8 and polyethylene. 9 Whereas the thermal treatment controls the type and degree of crystallinity, the type of fibre determines the ability to form a transcrystalline layer. The crystailinity of the fibre itself

Abstract Microcomposites of single-pitch-based carbon fibre reinforced J-Polymer are employed to investigate the mechanical role of the fibre matrix transcrystalline interphase. The transcrystalline interphuse in this semicrystalline thermoplastic system is varied by changing the crystallization kinetics, as determined by the thermal history. The kinetics of transcrystallization under isothermal conditions are presented as an example, and are used to form different transcrystalline interphase thicknesses. These affect the fibre fragmentation process that occurs while the specimens are cooled from the crystallization temperature to room temperature. This fragmentation process is attributed to residual thermal stresses, which can be calculated by assuming that it is controlled by Weibull statistics. Tensile loading of either longitudinal or transverse microcomposite specimens results in additional fragmentation, the extent of which is determined jointly by the thickness of the transcrystalline layer and by the yield strain of the matrix.

Keywords: pitch-based carbon fibre, J-Polymer, transcrystalline interphase, residual stresses, fibre fragmentation INTRODUCTION

The role of the fibre/matrix interphase in composite materials is currently the focus of an increasing number of studies. An interphase (interlayer) is a

Composites Science and Technology 0266-3538/93/$06.00 © 1993 Elsevier Science Publishers Ltd. 43

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Silvia lncardona et al.

is one of the factors contributing to the ability of the fibre to grow a transcrystalline layer on its surface. Hence, a transcrystalline layer can hardly grow spontaneously on the surface of a glass fibre, whereas the presence of a carbon fibre in a thermoplastic polymer may induce orientational behaviour even in the bulk polymer, as shown for a thermotropic main-chain aromatic co-polyester. ") Glass fibres can induce transcrystallization either by special surface treatment or as a result of a shear stress in the matrix due to fibre movement, 1~ or thermal stresses upon rapid cooling from the melt/z Regarding the ability of carbon fibres to grow a transcrystalline layer, a pitch-based carbon fibre, with its radial morphology, is more potent than a PAN-based carbon fibre with its circumferential morphology. This difference results from the fact that the radial morphology produces graphite basal plane edges at the fibre surface, which serve as nucleation sites for the transcrystalline layer. 7 Like the interphase in thermoset matrix composites, the transcrystalline layer is expected to have an effect on the mechanical performance of the composite. 6'9 It has been demonstrated, for example, that transcrystallinity enhances the fibre/matrix interfacial bond strength, 7 and increases the transverse strength of the composite. 6 Other effects of the transcrystailine layer, e.g. on the dynamic mechanical properties, 9 may be attributed to the claim in Ref. 13 that its elastic modulus is several times higher than that of the spherulitic material. The urgent need for additional scientific data regarding the role of the transcrystalline layer in different composite materials under a range of mechanical applications is combined in the present research with our ability to manufacture and test monofilament reinforced thermoplastic matrix composites. This ability was demonstrated recently in our study of transverse loading of monofilament reinforced microcomposites as a fragmentation technique for measuring the fibre compressive strength. 14 In the present work, carried out with microcomposites of high-modulus, pitch-based carbon fibre and J-1 Polymer, both longitudinal and transverse specimens were prepared and tested. Another study of the effect of thermal treatment on the type and kinetics of crystallization and transcrystallization,j5 which preceded this one, enabled different transcrystalline layers to be obtained by different thermal treatments.

EXPERIMENTAL The matrix of the monofilament microcomposites was J-1 Polymer, a polyamide homopolymer based on bis(para-amino cyclohexyl) methane (PACM), by Du Pont. The carbon fibre employed as monofilament reinforcements was a high-modulus, pitch-based fibre (PRD-172, Du Pont), of a nominal tensile strength

and modulus of 2-8 GPa j6 and 896 GPa, respectively. ~ The fibre had neither surface treatment nor size. Microcomposite samples approximately 10(I/~m thick were prepared by carefully sandwiching parallel individual monofilaments at 2cm intervals between four rectangles (8cm × 4cm) of dried, as received J-Polymer. The J-Polymer was pressed between two 1-mm-thick stainless steel plates of the same dimensions. Kapton polyimide sheet (Du Pont) coated with Freekote 44 and Freekote HMT release agents (Hysol) was placed between the J-Polymer and the steel plates to ensure easy release of the samples. Processing was performed at 320°C for 20min in a Carver press with minimal applied pressure, followed either by quenching in iced water, or by cooling at 2°C/min to the isothermal treatment temperature (270°C). After a measured period of isothermal treatment, the sample was quickly removed from the press and quenched in iced water. Miniature tensile samples 3 mm wide were carefully cut from the prepared sheets by a special cutting device under a microscope, ensuring that a single, straight monofilament either parallel or perpendicular to the sample axis could be observed half way along its length. Cardboard grips (1 cm 2) were glued to the ends of the sample with Poxipol room-temperaturecuring epoxy resin, giving a sample gauge length of 20 mm. Postcuring was performed at 60°C for 1 h under pressure to prevent sample slippage during testing. Both longitudinal and transverse samples were loaded at a rate of extension of 0.83/~m/s in an elaborate video-monitored tensile testing system described in detail elsewhere. 18 The lengths of the fibre fragments at each stress level were measured by means of a Colorado video micrometer (Model 305A), with an estimated accuracy of 5/~m. Hot-stage microscopy was carried out with circular specimens, 5 mm in diameter, punched out of the quenched microcomposite sheet. The specimen was remelted on the hot stage at 300°C, and then cooled rapidly to the treatment temperature. The equipment comprised a Nikon stereoscope, equipped with a Mettler FP82 hot stage/FP80 central processor. The resolution of the thickness measurement of the transcrystalline layer was 2.5 jum.

RESULTS AND DISCUSSION Isothermal transcrystailization The transcrystalline layers discussed here were grown isothermally from the melt (a full account of the thermal and kinetic results is given elsewherelS), and their size and shape depended on the isothermal crystallization temperature and time. For a qualitative demonstration, Fig. 1 presents a series of photomicrographs for a range of crystallization temperatures. It is seen that different thicknesses and perhaps

d)

ic)

Fig. 1. Photomicrographs of isothermal crystallization and transcrystallization at different temperatures: (a) 275°C, (b) 265°C, (c) 255°C and (d) 245°C.

[]

[]

~°

46

Silvia Incardona et al. 60"

E

crystallization, where it results from the opposing effects of increasing the driving force of crystallization and decreasing the molecular movement with decreasing temperature. 19 This indicates that the kinetics of transcrystallization are similar to those of bulk crystallization.

40"

30

z

20

=~ El

10" 0~

~

100

200

time ( mln )

Fig. 2. The isothermal growth with time of the thickness of the transcrystalline interphase for three temperatures: ([~)

275°C, (0) 265°C and (ll) 260°C.

250'

E

150

50

-50 230

, 240

, 250

, 260 Temperature (°C)

' 270

280

Fig. 3. The isothermal growth rate of the transcrystalline interphase as a function of temperature. different crystallization habits are obtained for different temperatures. The isothermal rate of growth of the transcrystalline layer was linear, as shown for three temperatures in Fig. 2. The rate of growth exhibits a maximum at 240°C, shown in Fig. 3. The presence of a maximum is known for spherulite

Residual thermal stresses Every study of the mechanical performance of semicrystalline thermoplastic matrix composites ought to consider the build-up of residual thermal stresses in these materials. The residual thermal stresses, which result from differential thermal shrinkage of the consitutents on cooling during processing, are expected to be highest for semicrystalline thermoplastic matrices, e° This is attributed to a large amount of matrix shrinkage that occurs during bulk crystallization, which in turn depends on the thermal conditions. For specific combinations of thermal conditions and constituent material properties, the residual thermal stresses may cause permanent damage, such as the observed matrix cracks in Kevlar-49 aramid fibre reinforced J-Polymer composites.21.22 When the residual thermal stresses are sufficiently high, they can result in compressive failure of the carbon fibre. This, in particular, is feasible for high-modulus, pitch-based carbon fibres, whose compressive strength is very low. z3 lndeed, fibre compressive failure occurred repeatedly during the manufacture of the microcomposites in the present study, where a high-modulus, pitch-based carbon fibre was used. Figure 4 shows a typical compressive break distinguished by its diagonal path. In fact, depending on the thermal conditions of crystallization, a fibre fragmentation process in compression took place,

Fig. 4. A typical fibre compressive break, resulting from a build-up of residual thermal stresses as the microcomposite is cooled from the crystallization temperature to room temperature.

Transcrystalline interphase in J-Polymer microcomposites 12"

E

o

~E

Table 1. Calculated residual thermal stresses at different transcrystalUne interphme thicknesses, obtained isothermally at 270°C

m

E

47

10"

6'

Duration of treatment (min)

Thickness of interphase (/~m)

15 30 60

11.7 23-4 46-8

Fragment Compressive length (/~m) strength of fibre (GPa)

4. 0

20

30

40

5'0

60

70

time ( rain )

Fig. 5. The break density in the compressive fragmentation process induced by the residual thermal stresses after different isothermal crystallization times at 270°C. resulting in a series of compressive fibre breaks. This resulted from the residual thermal stresses built up during the manufacture of the microcomposite. It was soon noticed that the fragmentation process during the manufacturing stage depended on the thermal conditions. The most obvious dependent variable was the number of compressive breaks per unit fibre length. Figure 5 presents the fragmentation results obtained from microcomposites prepared isothermally at 270°C, followed by iced water quenching. Each point represents an average of at least seven specimens. It is seen that the number of breaks per unit length decreases in an approximately linear fashion, as the time of the isothermal crystallization treatment increases. Fragment lengths of around 100, 160 and 280/~m are obtained for 15, 30 and 60 rain isothermal treatments. This behaviour is associated with the thickness of the transcrystalline layer, which was shown to grow linearly with time (see Fig. 2). It is also noted that these fragment lengths correspond with the reported range of critical lengths of various carbon fibres in different matrices of 500/,m or less. 24 These observations suggest that a better developed transcrystalline layer produces lower residual thermal stresses, It appears that a thick transcrystalline layer produces some sort of protective action by relieving the residual thermal stresses that originate in the differential thermal expansivity of the matrix and the fibres. Thus, when a thin transcrystalline layer exists, the differential thermal expansivity over a temperature span of 270°C (based on isothermal treatment at 270°C, followed by iced water quenching) induces a residual stress sufficiently high to result in the ultimate compressive fragmentation, namely, of the order of the fibre critical length. The ability of a thick transcrystalline layer to relieve partially the thermal stresses may stem from its expected high elastic modulus in the fibre direction, which is perhaps several times higher than that of the spherulitic material in the bulk matrix. 9'~3 This implies that, with a thick interlayer, a lower strain is experienced by the fibre/transcrystaUine layer entity, inducing smaller stresses in the fibre.

100 160 280

2.06 1.99 1.90

As the extent of fragmentation reflects the residual stress level in the fibre, which in turn, for a given temperature span, depends on the thickness of the transcrystalline layer, it is suggested that the fragmentation process be used to assess residual thermal stresses. The quantitative approach is based on a recent paper 14 demonstrating how the entire compressive loading history of a single fibre in a fragmentation test could be utilized to measure the size effect on fibre strength. It was shown that, as in a tensile loading situation, 25 the size effect in a compressive loading situation could be determined by a continuously monitored experiment which yields a set of fragment lengths as a function of the applied stress, according to eqn (1), as derived from the Weibull distribution: S a = ¢~flGfu/$[I~(1 + 1 / ~ ) ] fl

(1)

where t~ and fl are, respectively, the Weibull scale and shape parameters for strength, F is the gamma function and of, is the compressive strength of the fibre for a fragment length sa (which becomes la at break saturation). Assuming that a[~ is equal to the residual thermal stress in the fibre, eqn (1) can be used, provided that the WeibuU scale and shape parameters are known for the specific fibre. From the experimental results in Fig. 5, the residual stress in the fibre can be calculated for different transcrystalline layers, as presented in Table 1. The WeibuU scale and shape parameters used in the calculation are a~ = 1.0 GPa and fl = 12.0, assumed for high-modulus, pitch-based carbon fibre. Although these values of the scale and shape parameters are taken as an example to demonstrate how eqn (1) can be used to calculate the residual stresses, they are probably realistic for the high-modulus, pitch-based carbon fibre in compression. A scale parameter of 1.0 GPa is reasonable in view of a value of 0.163% reported for the compressive failure strain of a 662GPa modulus pitch-based fibre, 26 while the nominal modulus of the fibre used in this study was 896 GPa. 17 Tensile and compressive fragmentation Each of the above described microcomposites was tested by tensile stressing both longitudinally and transversely to the fibre direction. Loading was

Silvia lncardona et al.

48

(a)

@

J

i !

(b)

Fig. 6. Matrix cracking (a) and matrix yielding (b) in longitudinal microcomposlte specimens. continued until a point of matrix yielding or cracking was identified as shown in Fig. 6. This occurred at about 2 - 3 % strain (see Fig. 8). Thereafter, the specimens were re-examined in order to record additional fibre breaks. In the longitudinal monofilament microcomposites the additional fibre breaks

resulted from a tensile fragmentation process due to the external loading, as analysed recently. 25 In the transverse microcomposites they resulted from a recently described compressive fragmentation process due to a Poisson transverse contraction. ~4 Figure 7 compares the pre- and post-stressing fragmentation

Transcrystalline interphase in J-Polymer microcomposites

49

12 E E m

"6

60 10 8" n 40 E

II

6'

44

sl 2O

2 0

2 0'

3'0

"

4 0'

50 ' ' " 6 0' time ( rain )

70

Fig. 7. Fibre break densities in microcomposites after different isothermal times at 270°C: ([~) induced by residual thermal stresses only, (11) additional tensile stress induced fragmentation in a longitudinal specimen, and (~l~) additional compressive stress induced fragmentation in a transverse specimen. results. It is seen that in the 15min isothermal treatment the post-loading fragment density was equal, within the experimental scatter, to that of the pre-loading. This observation substantiates the assumption that in these conditions an ultimate fragmentation density, corresponding to fragment lengths of the order of or below the fibre critical length in tension, has already occurred as a result of thermal stresses. For the other treatments, mechanical stressing of the microcomposites increased the average number of breaks per unit length. However, the ultimate fragmentation density could not be reached, probably because the external stress level was limited by the tensile yield strain of the matrix. The observation that the extent of the external stressinduced fragmentation process was higher in the transverse microcomposite, compared with the longitudinal, can be explained by the fact that the fibre was already under a compressive stress state, and also by its lower compressive strength compared with the tensile strength. It has been argued that an exhaustive fragmentation test requires a matrix yield strain that is at least three times higher than that of the fibre. 27 The tensile yield strain and stress of the matrix could be detected from the stress/strain curve, as exemplified in Fig. 8. The yield stress (at 0.01 offset) ranged, according to the thermal treatment, from 25-0 MPa in the quenched material to 47.3 MPa in the maximum crystalline material, while the yield strain was constant at approximately 0.035. The empirical observations of matrix cracking and yielding, in fact, sustain the claim that the fragmentation processes in the 30 and 60min treatments are not exhausted, and the observed average fragment lengths are longer than the fragment length at break saturation. Although an alternative possibility, that a different critical length exists for each transcrystalline thickness owing to different interracial interactions, cannot be ruled out, it seems less probable. The different thermal treatments are expected to result in different transcrystalline

0

0.01

0.02

0-03

I 1 0.04 0.05 0.06 Strain ( m m l m r n )

f 0.07

I 0.08

I 0.09

(a) 60

g~ 40 ~ 20!

0

0-01

002

0.03

0 . 0 4 0 . 0 5 0-06 Strain ( r n m / m r n )

L

007

0-08

I

0@9

(b) l~g. 8. Stress/strain traces of longitudinal or transverse microcomposite specimens: (a) isothermally treated at 270~C for 60 rain; (b) quenched. thicknesses and not to affect the specific fibre]matrix interaction at the surface of the fibre. For example, a recent study on poly(phenylene-sulphide)/carbonfibre composites emphasizes the importance of the number of nucleation sites on the fibre surface (and in the matrix). 28 As this number is expected to be constant regardless of the thermal conditions, so is the bond strength, and the argument that the critical length is independent of the thickness of the transcrystalline interphase is valid. CONCLUSIONS An extensive study of the kinetics of transcrystalline growth, whose partial results are described in this paper, reveals a number of aspects of the mechanical and physical significance of transcrystallinity. It is obvious that the thickness of the transcrystalline layer affects the level of residual thermal stresses, as manifested by the average fragment length, in the composite. A thicker layer, of probably a higher elastic modulus, results in lower thermal stresses, and hence in larger fragment lengths. Assuming that each fragment length is related to a single compressive strength value by the WeibuU distribution, and that this fibre strength is identical to the residual thermal stress, a method for the determination of the residual

50

Silvia lncardona et al.

stress is proposed, which is based on a knowledge of the Weibull scale and shape parameters. Further compressive or tensile fragmentation under the respective loading modes, and in the absence of matrix yielding, may lead to a determination of the fibre critical length, which is probably independent of the thickness of the transcrystalline layer.

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London, 1991 p. 111. 13. Folkes, M. & Hardvich, S., J. Mater. Sci. Lett.. 6 (1987) 656. 14. Wagner, H. D., Gilbert, A. H., Migliaresi, C. & Marom, G., J. Mater. Sci., 27 (1992) 4175. 15. lncardona, S., Fambri, L., DiMaggio, R., Marom, G. & Migliaresi, C., Crystallization and transcrystallization in J-Polymer and its carbon fibre composites. J. Mater. Sci., (in press). 16. Wagner, H. D., Aronhime, J. & Marom, G., Proc. R. Soc. Lond., A428 (1990) 493. 17. Gilbert, A. H., Goldstein, B. & Marom, G., Composites, 21 (1990) 40818. 18. Wagner, H. D. & Steenbakkers, L. W., J. Mater. Sci., 24 (1989) 3956. 19. Sperling, L. H., Introduction to Physical Polymer Science. Wiley, New York, 1986. 20. Nairn, J. A. & Zoller, P. J. Mater. Sci., 20 (1985) 355. 21. Lee, W. J., Fukai, B. K., Seferis, J. & Chang, I. Y., Polym. Comp., 6 (1988) 473. 22. Marom, G. & Chen, E. J-H., Composites Sci. Tech., 29 (1987) 161. 23. Ohsawa, T., Miwa, M., Kawade, M. & Tsushima, E., J. Appl. Polym. Sci., 39 (1990) 1773. 24. Asioun, El. M., Donnet, J. B., Guilpain, G., Nardin, M, & Schultz, J., J. Mater. Sci., 24 (1989) 3504. 25. Yavin, B., Gallis, H. E., Scherf, J., Eitan, A. & Wagner, H. D., Polym. Comp., 12 (1991) 436. 26. Prandy, J. M. & Hahn, H. T., SAMPE Quarterly, 22(2) (1991) 47. 27. Verpoest, I. & Desaeger, M., In Interracial Phenomena in Composite Materials '91, ed. I. Verpoest & F. Jones. Butterworth-Heinemann, London, 1991, p. 9. 28. Caramaro, I., Boudet, A., Chabert, B., Chauchard, J. & Vu-Khanh, T. In Interracial Phenomena in Composite Materials '91, eds I. Verpoest & F. Jones. Butterworth Heinemann, London, 1991, p. 251.