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PEEK composites. Part 1. Crystallinity and interface adhesion
Composites: Part A 31 (2000) 517–530 www.elsevier.com/locate/compositesa
Cooling rate influences in carbon fibre/PEEK composites. Part 1. Crystallinity and interface adhesion Shang-Lin Gao, Jang-Kyo Kim* Department of Mechanical Engineering, Hong Kong University of Science and Technology, Clear Water Bay, Hong Kong, People’s Republic of China Received 17 June 1999; received in revised form 22 November 1999; accepted 19 January 2000
Abstract The effect of cooling rate on the fibre–matrix interface adhesion for a carbon fibre/semicrystalline polyetheretherketone (PEEK) composite was characterised based on the fibre fragmentation, fibre pullout and short beam shear tests. The interface adhesion was correlated to the degree of crystallinity and the crystalline morphology, as well as the bulk mechanical properties of neat PEEK resin, all of which were in turn controlled by cooling rate. It was shown that the interface bond strength decreased with increasing cooling rate; the tensile strength and elastic modulus of PEEK resin decreased, while the ductility increased with increasing cooling rate through its dominant effect on crystallinity and spherullite size. The improvement of crystalline perfection and flattened lamella chains with high crystallinity at the interphase region were mainly responsible for the strong interface bond in composites processed at a low cooling rate. The interphase failure was characterised by brittle debonding in slow-cooled composites, whereas the amorphous PEEK-rich interphase introduced in fast cooled specimens failed in a ductile manner with extensive plastic yielding. 䉷 2000 Elsevier Science Ltd. All rights reserved. Keywords: Cooling rate; Strength; B. Mechanical properties
1. Introduction Thermoplastic matrix composites have been studied extensively in an effort to best utilise the potential advantages of high fracture toughness, high temperature resistance, repairability and ease of manufacture. It has long been recognised that the mechanical properties of semicrystalline polymers and the composites made therefrom are dependent on the crystallinity and crystalline morphology, which are strongly affected by the processing conditions [1,2]. The presence of carbon fibres within the matrix induces nucleation and growth of crystallites perpendicular to the fibre surface, i.e. transcrystallisation, which has a considerable influence on the fibre/matrix interface interaction and the failure behaviour in both the matrix and the interphase region [3]. Optimisation of the bulk mechanical properties of thermoplastic matrix composites, therefore, requires better understanding of the effect of processing conditions on the interphase properties. The interactions at the interphase region in semicrystalline thermoplastic composites depend on a number of
* Corresponding author. Tel.: ⫹ 852-2358-7207; fax: ⫹ 852-23581543. E-mail address: [email protected] (J.-K. Kim).
factors, such as matrix morphology, fibre surface condition, presence of residual stresses, moduli of the fibre and matrix, as well as the presence of reactive functionalities [4]. The majority of these characteristics are determined by processing conditions, including moulding temperature, cooling rate, holding time/temperature and annealing conditions. A great deal of studies have been directed toward understanding the influences of these parameters, but the results reported hitherto displayed no apparent consensus. For example on the effect of cooling rate, an increase in the interfacial shear bond strength (IFSS) was noted with increasing cooling rate for glass fibre/polypropylene (PP) [5] and carbon fibre/polycarbonate (PC) systems [6]. This result was attributed to the improved wetting behaviour, high ductility, fracture energy and low residual stresses in the matrix and the inhibition of molecular weight segregation or weak boundary layer that arose from the high cooling rate. However, there were contradictory reports for carbon fibre/PP matrix [7], carbon fibre/polyphenylene sulfide (PPS) [8], carbon fibre/polyethylene terephthalate (PET) [9] and alumina fibre/PP systems [10], showing a totally opposite trend, or no significant cooling rate dependence. The high IFSS with decreasing cooling rate was partly due to the formation of transcrystalline interface and partly to the improved adsorption capability.
1359-835X/00/$ - see front matter 䉷 2000 Elsevier Science Ltd. All rights reserved. PII: S1359-835 X( 00)00 009-9
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Table 1 Reported interface shear strength of carbon fibre/PEEK composites System
IFSS (MPa)
Testing methods
Processing condition
Reference
AS4/PEEK AS4/PEEK AS4/PEEK AS4/PEEK
38–43 50–97 70–100 “Good interface adhesion for slow cooling”
Pullout Fragmentation Pullout SEM observation
Isothermal crystallisation Cooling rate: ⬃400⬚C/min Quenching and annealing Cooling rate: 2–63⬚C/min
[9] [13] [14] [15]
The effects of transcrystallinity on the IFSS and the bulk mechanical properties of composite were also found to be inconsistent. It was argued that large thermal stresses might build up during the transcrystallisation process, which in turn reduced the IFSS, whereas the increasing degree of crystallinity of matrix material effectively decreased the mismatch in moduli between the fibre and matrix, and thus enhancing the stress transfer efficiency across the interface. The formation of such transcrystalline layer, thus, improved the tensile modulus, strength and fracture toughness of short carbon fibre/PEEK matrix composites [3]. In contrast, Marom [11] attributed the effect of transcrystalline layer to a preferred crystallite orientation relative to the fibre, thereby providing the surrounding matrix with a high rigidity and reduced thermal expansion in the fibre direction, which effectively lowered the residual thermal stresses. Meanwhile, transcrystallinity was found to have conflicting influences on the IFSS, depending on the matrix material [12]. For carbon fibre/PEEK systems, the presence of transcrystalline interphase layer gave rise to the IFSS, whereas for carbon/PES system, the reverse was true, although the differences were only marginal in both cases. Even for an identical AS4 carbon fibre/PEEK system with which the present study is mainly concerned, the measured IFSS varied widely from 40 to 110 MPa, due mainly to the largely different processing conditions [12–15], as presented in Table 1. The use of different testing techniques and different data reduction methods were also partly responsible for the large data variation. The inconsistency on the effects of cooling rate and crystallinity warrants a systematic study to clarify the relationship between cooling rate, morphology and the IFSS. The present study is the continuation of our previous work [16–18] on the improvement of mechanical performance and damage tolerance of carbon fibre/ PEEK matrix composites through the optimisation of processing conditions and control of the fibre–matrix interphase properties. The dependence of IFSS on matrix crystallinity was evaluated over a wide range of cooling rate from 1 to 2000⬚C/min, using a variety of thermomechanical characterisation techniques and mechanical testing methods. Model microcomposites were employed to measure the IFSS in the single fibre pullout and fragmentation tests.
2. Experimental 2.1. Materials and characterisation of crystallinity All specimens used in the present study were prepared from Hercules AS4 carbon fibres, PEEK powder (Victrex) and APC-2 prepregs containing continuous AS4 carbon fibre of 61% by volume. The PEEK powder was a similar material to the matrix of APC-2 prepregs according to the manufacturer (ICI Fiberite). The crystallinity of both neat PEEK resin and carbon fibre/PEEK composites was measured using a differential scanning calorimeter (DSC, SETAPRAM-92). Specimens of approximately 10 mg were analysed in nitrogen at a heating rate of 20⬚C/min. The following equation was used to calculate the degree of crystallinity, X: X
DHm ⫺ DHc DHf 1 ⫺ a
1
where DHm and DHc are the enthalpy of fusion at melting point as measured by the area under the endothermic peak, and the enthalpy of crystallisation as measured by the area under the exothermic crystallisation peak, respectively. DHf is the enthalpy of fusion of fully crystalline PEEK, which is taken 130 J/g [19]; and a is the mass fraction of carbon fibre in the composite. An optical microscope equipped with polarised light filters was used to characterise the crystalline morphology of the interphase and the surrounding matrix. The surface of metallographically polished composite specimens was etched using a mixture of 1 wt% potassium permanganate, surfuric acid (five parts), ortho-phosphoric acid (two parts) and distilled water (two parts by volume). 2.2. Single fibre pullout and fragmentation tests The fibre pullout and fibre fragmentation tests are amongst the most popular techniques to measure the fibre–matrix interface bond quality [20]. Along with the usual difficulties in measuring the interfacial shear strength based on single fibre model composites because of the small diameter of carbon fibres, the PEEK matrix, in particular, presented problems in specimen fabrication due to the high melt viscosity and high melting temperature. These problems necessitated the development of special methods to fabricate single fibre composite specimens at different
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different thickness depending on the desired cooling rate, which were then cooled between two aluminium plates. Cooling rates above 1000⬚C/min were achieved by directly quenching the moulds in a water bath of ambient temperature. The moulds were wrapped using aluminium foil to avoid direct contact with water. The cooling rate was monitored using a Datapag tracker with thermocouples embedded in the specimen. The pullout and fragmentation experiments were performed on a Minimat testing machine at a cross-head speed of 0.01 mm/min and 10 mm/min, respectively. The fibre/matrix interface morphology and the fibre fragment lengths were simultaneously monitored using a polarised light microscope, a video recorder and an image analyser. The whole testing set-up is schematically shown in Fig. 2. The critical transfer length, lc, was calculated from the measured mean fibre fragment lengths, l [21]: lc 4l=3
Fig. 1. Schematic of specimen preparation for: (a) single-fibre pullout; and (b) fibre fragmentation tests.
cooling rates. Specimens were prepared by partially or wholly embedding a carbon fibre within the PEEK powder, which was contained in a mould, as shown schematically in Fig. 1. The moulds were heated to 400⬚C in an oven for 10 min, which were then allowed to cool as detailed in the following. Cooling rates below 10⬚C/min were achieved by cooling the moulds in a preheated oven of different initial temperatures. Moulds for cooling rates between 10 and 1000⬚C/min were shielded using polyimide films of
2
The critical aspect ratio, lc/d and the modulus ratio, Ef/Em, of fibre to PEEK matrix were calculated for the AS4 carbon fibre of an average diameter d 7:7 mm and modulus Ef 234:4 GPa [22]. In view of the fact that the maximum loads, F, measured from the fibre pullout tests consistently displayed a linear dependence on the embedded fibre length, L, the IFSS was calculated using the simple linear relation: IFSS
F pdL
3
The baseline mechanical properties of neat PEEK resin were measured in uniaxial tension of PEEK films according to the specification ASTM D11708. The interlaminar shear strengths were also measured from the short beam shear (SBS) tests of unidirectional carbon fibre/PEEK
Fig. 2. Schematic of Minimat testing fixture placed on a microscope stage.
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Fig. 3. DSC thermograms for: (a) neat PEEK resin; and (b) carbon fibre/PEEK composites processed at different cooling rates.
composite laminates that were fabricated from APC-2 prepregs and processed at different cooling rates. The SBS tests were conducted on a MTS 858 universal testing machine at a cross-head speed of 1.3 mm/min, according to the specification ASTM D2344. The fracture surfaces of the composites were examined using a scanning electron microscope (Cambridge Stereoscan 200).
3. Results and discussion 3.1. Crystallinity and morphology The crystallinity contents were measured using DSC and the corresponding DSC thermograms are shown in Fig. 3 for the neat PEEK resin and carbon fibre/PEEK composites
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Table 2 Effect of cooling rate on crystallinity of neat PEEK resin and carbon fibre/PEEK composite Cooling rate (⬚C/min) 1 70 160 600 1000 1500 a b
Tc (⬚C) – – – 188 (182) 187 188 (183)
DHc (J/g)
Tm (⬚C) 351 (346) 351 (344) 352 352 (342) 352 351 (342)
a
– – – 3 具3典 13 14 具15典
DHm (J/g) 49 38 37 37 38 36
具42典 具36典 具36典 具31典
b
Crystallinity (%) 38 具33典 30 具28典 28 26 具25典 19 17 具12典
Values in parentheses ( ) represent those for carbon fibre/PEEK composite. Values in parentheses 具 典 represent those for carbon fibre/PEEK composite taking into account the matrix weight fraction of 32%.
processed at different cooling rates. The PEEK resin and the composite processed at a very high cooling rate, say at 1000⬚C/min or above, showed only a marginal difference between the areas of crystallinity peak and melting peak, indicating a very low degree of crystallinity. The semicrystalline PEEK resin and the composite processed at a low cooling rate displayed almost a negligible crystallisation peak, giving rise to a high degree of crystallinity. The results are summarised in Table 2 and the crystallinity is plotted as a function of cooling rate in Fig. 4. Also included in this figure is the compilation of crystallinity values reported in the literature [23–29]. The data indicates considerable scattering, in part attributable to the different techniques employed to measure crystallinity, such as DSC, wideangle X-ray spectroscopy (WAXS) and the density method [30], as well as to the different types of fibres and matrix materials studied. Even so, the general trend clearly indicates that the higher the cooling rate, the lower the crystallinity. For both the neat PEEK resin and carbon fibre/PEEK composites, the degree of crystallinity decreased as the cooling rate increased (Fig. 4). It is well known that
crystallinity, melting point and cooling rate are closely interrelated. The melting peak temperature of a semicrystalline polymer composite is determined by several factors, such as crystal size, fold surface energy and plate thickness [31]. The size of sphrullite in turn is strongly influenced by cooling rate: the lower the cooling rate, the larger the spherullite size with a higher degree of molecular perfection and a lower melting temperature [27,32,33]. The correlation between the spherullite size and cooling rate is clearly seen from the cross-polarised optical micrographs given in Fig. 5. The spherullites in the PEEK resin obtained after very slow cooling, say 1⬚C/min, was as large as a few hundred mm in diameter, while they were too small to be detected using an optical microscope once a very high cooling rate was used, say 2000⬚C/min. The crystallite size dependence on cooling rate can be explained from the viewpoint of molecular movement: the mobility of polymer chains is discouraged at a high cooling rate, thus limiting the ability of the chains to diffuse into the growing crystal front. Another significant implication of Fig. 5 is that the fibre surface influenced the nucleation and growth of spherullites within the matrix, particularly at a low cooling rate. It is
Fig. 4. Degree of crystallinity as a function of cooling rate for PEEK resin and carbon fibre/PEEK composites. Measurement methods are shown in ( ).
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Fig. 6. (a) Scanning electron micrograph of fibre-nucleated spehrullites on an etched cross-section of a slowly cooled carbon fibre/PEEK composite. (b) Schematic diagram showing the growth of highly oriented spherullites from the fibre surface [36]. Fig. 5. Optical micrographs of PEEK matrix morphologies around isolated carbon fibres at cooling rates: (a) 1⬚C/min; (b) 200⬚C/min; (c) 1000⬚C/min; and (d) ⬃2000⬚C/min.
well known that nucleation and growth of spherullites from the surface of high modulus carbon fibres is often related to the graphitic nature of the fibre surface rather than the fibre surface chemistry [34,35]. A further proof for these fibre surface-induced spherullites is presented in Fig. 6. The spherullites emerging from the fibre surface impinged each other within the narrow matrix bed due to the presence of rigid fibres of high volume fraction. The schematic illustrated in Fig. 6(b) [36] depicts spherullites of a high order of alignment and compact packing. It is believed that the crystalline structure developed at a slow cooling rate is better defined and larger in size with more distinct and thicker lamellae than that developed at a high cooling rate. The fast-cooled composite generally displayed a mixture of isolated and fibre-nucleated spehrullites, whereas
Fig. 7. Tensile stress–strain curves of neat PEEK resin processed at different cooling rates.
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Fig. 8. Ductility, tensile strength and modulus of neat PEEK as a function of cooling rate.
the slow-cooled composite only revealed much larger fibre nucleated spherullites [1]. It is interesting to note that both the degree of crystallinity and the melting peak temperature of the composite were slightly lower than those of the neat PEEK resin for the same processing condition (Table 2). A similarly reduction in melting temperature for a carbon/PPS composite [37] and in crystallinity for a carbon/PEEK composite [38] were previously reported. It seems that the presence of densely packed fibres within the matrix severely suppressed the growth of spherullites, effectively offsetting the aforementioned ameliorating effect of fibre surface acting as preferred crystallite nucleation sites. 3.2. Tensile properties of matrix Typical engineering tensile stress-strain curves of neat PEEK resin are presented in Fig. 7, and the average values of the major mechanical properties are summarised in Fig. 8. It is obvious that both the elastic modulus and the strength of PEEK increased, while the ductility decreased with
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decreasing cooling rate. As would be expected from the DSC results (Fig. 4), the high crystallinity with ordered crystallites obtained at a low cooling rate gave rise to a high stiffness of the material, causing yielding to be more difficult with dramatically reduced ductility. A high stiffness and high strength means that it is more difficult to induce slips in the crystal blocks due to the higher orderliness and thicker lamellae. Fig. 7 further indicates that there were distinct stress drops after yield point in the specimens cooled at a high cooling rate, say at or above 1000⬚C/min, while the stress drop was totally absent for the specimens processed at low cooling rates. Similar tensile behaviour for the annealed PEEK resin suggests [39] that the stress drop at the yield point was hidden by the occurrence of postyield hardening behaviour due mainly to the improvement of crystalline perfection and lamella thickening taking place during the annealing treatment. It is expected that a slow cooling process promotes the formation of more perfect crystallites with a higher proportion of thick lamellae than a fast cooling process. It is also worth noting that an increase of cooling rate from 1⬚C/min to about 600⬚C/min resulted in a significant change in all properties measured, including the ductility, tensile strength and modulus. Further increase in cooling rate beyond 600⬚C/min did not much vary these values. This distinct behaviour may be a direct reflection of the change in crystallinity, and similar trends with respect to cooling rate were noted for the IFSS and the composite interlaminar shear strength (ILSS), as discussed in Section 3.3.2. 3.3. Interaction at the interphase region 3.3.1. Fibre fragmentation test There was considerable scattering in the fibre fragment length, l. Thus, the statistical variation of l was described by the cumulative probability, P, which was calculated from: P l
n l N ⫹ 1
4
where n(l) is the number of fibres whose lengths are equal to or shorter than l, and N is the total number of fragmented fibres at saturation. The experimental data were analysed using the simple cumulative probability theory proposed previously [40]: P l 1 ⫺ exp⫺ l=l0 m
Fig. 9. Cumulative fibre fragment length for different cooling rates.
5
where m and l0 are the Weibull modulus (or shape factor) and the characteristic length (or scale factor) of fragmented fibres, respectively. A strong correlation was noted between the fibre fragment length distribution and cooling rate. The cumulative distributions are presented in Fig. 9, and the Weibull parameters as well as the mean fibre fragment lengths are summarised in Table 3. It is obvious from these results that a shift to the shorter fragment lengths corresponded to a slower cooling rate. Meanwhile, the
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Table 3 Weibull parameters for fibre fragmentation lengths at saturation Cooling rate (⬚C/min)
Weibull modulus, m
Weibull scale parameter lo (mm)
Mean fibre fragmentation length ^ SD (mm)
490 600 1000 1500 1800
2.88 2.81 3.55 3.16 2.91
399 487 525 550 610
350 ^ 45 432 ^ 82 470 ^ 110 491 ^ 64 543 ^ 114
Weibull modulus, m, varied rather irregularly with the cooling rate without showing any general trend, indicating that the distribution of fibre fragment lengths was insensitive to cooling rate. The critical fibre aspect ratio, lc/d, calculated from the mean fragment length at saturation is plotted as a function of cooling rate in Fig. 10. There is approximately a linear relation between these parameters: the higher the cooling rate, the higher the critical aspect ratio. This indicates that the efficiency of stress transfer across the fibre–matrix interface is inversely proportional to cooling rate. The high stress transfer efficiency at a low cooling rate may be explained in terms of improved mechanical characteristics of the matrix material and, to a certain extent, the high interface bond quality. In fact, the major factors behind the variation of stress transfer efficiency measured by fibre fragment lengths are rather matrix-related: the elastic modulus of the matrix material played a dominant role. The critical aspect ratio, lc/d, is plotted as a function of the modulus ratio of fibre to matrix, Ef/Em, in Fig. 11, clearly indicating there is a strong correlation between these parameters. This again confirms that the degree of crystallinity and crystalline morphology were mainly responsible for the variation of stress transfer efficiency through the effect of cooling rate. The above observation is quite consistent with experimental evidence on other composite systems that lc/d varied approximately
with (Ef/Em) 1/2 [41–43], as the early shear lag model suggests [44]. It should be mentioned here that Figs. 9 and 10 contain only the data corresponding to the specimens that were processed at high cooling rates and were loaded until saturation of fibre fragmentation. The specimens processed at cooling rates lower than about 400⬚C/min, however, tended to fail prematurely due to the cracks originating from the broken fibre ends, disallowing the measurement of saturation fibre fragment lengths. Also associated with this observation is the cooling rate-dependent failure behaviour of the fibre–matrix interface near the broken fibre ends. Fig. 12 presents typical microphotographs showing distinct failure modes. A transverse matrix crack in a peculiar diamondshape is clearly seen for the specimen with a stiff matrix and/or a strong interface, corresponding to a low cooling rate. Similar transverse cracks appeared at almost all broken fibre ends. In contrast, a gap between two broken fibres is detected for the specimens with a compliant, ductile matrix and/or a weak interface, corresponding to a high cooling rate. With increasing cooling rate, the failure mode changed from matrix crack propagation to interfacial debonding, with the empty cylindrical channel representing a debonded interface [45,46]. The failure mode change is a direct reflection of the improvement in ductility and fracture resistance of the matrix.
Fig. 10. Critical fibre aspect ratio, lc/d, versus cooling rate for carbon fibre/PEEK composites.
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Fig. 11. Critical fibre aspect ratio, lc/d, versus modulus ratio of fibre to matrix, Ef/Em for carbon fibre/PEEK composites.
3.3.2. Fibre pullout test The results obtained from the fibre pullout test complement and, most often, support the implications derived from the fibre fragmentation test. The interface failure mechanisms taking place during the fibre pullout processes were studied based on the sequential polarised microphotographs and the typical pullout load–displacement records, as shown in Fig. 13. The interface debonding process in the specimen processed at a very low cooling rate, e.g. 1⬚C/min (Fig. 13(a)), was characterised by brittle fracture with little matrix deformation along the fibre. As a consequence, almost a linear increase in pullout load was noted until it reached the maximum. In the fast cooled specimens at 1800⬚C/min (Fig. 13(b)), however, the rising part of the load–displacement trace exhibited an abrupt change in slope, indicating a rather stable interphase failure process between the two peaks. A typical failure mechanism was ductile yielding of the interphase, whose details are discussed in the
Fig. 12. Photomicrographs of stretched carbon fibre/PEEK fragmentation specimens cooled at: (a) 500⬚C/min; and (b) 1800⬚C/min.
following. At stages 1–3, the interphase deformed viscoelastically until yielding took place, as indicated by the bright fringe patterns around the fibre surface. At stages 3–5, interface debonding propagated stably from the embedded fibre end along with extensive plastic yielding of the interphase. At stages 5–6, the interface crack propagated unstably, and the fibre was completely pulled out. A thick layer of permanently deformed material is clearly seen surrounding the fibre socket after the complete pullout (see the photograph for stage 6). It appears that yielding took place at the early stage of pullout, and the thickness of the matrix layer affected by plastic deformation was approximately twice the fibre diameter for a typical specimen, as shown in Fig. 14. Examination of the fracture surfaces of pulled out fibres revealed that: for the slow cooled specimens, the fibre was relatively clean and free of matrix residue, indicating rather brittle interfacial debonding. In contrast, for the fast cooled specimens, some residual pieces of PEEK resin were seen adhering to the fibre surface, indicating plastic deformation of the soft, ductile matrix before interphase failure. Having studied the interphase phenomena during the fibre pullout test, which were shown to be very sensitive to cooling rate, the maximum pullout-force data were analysed in the following. The maximum pullout-force taken from the load–displacement curves were found to increase linearly with embedded fibre length for all specimens processed at different cooling rates, as shown in Fig. 15. The IFSS values calculated based on Eq. (3) are presented in Fig. 16 along with the results from the short beam shear test and DSC analysis. An increase in cooling rate from 1⬚C/min to about 600⬚C/min resulted in a significant reduction in all properties measured, including IFSS, ILSS and the degree of crystallinity. Further increase of cooling rate beyond 600⬚C/min, however, did not much vary the IFSS and ILSS, while the degree of crystallinity decreased further with cooling rate. Comparison of Figs. 8 and 16(a) revealed that the dependence of IFSS and ILSS on cooling rate was
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Fig. 13. Sequential microphotographs showing fibre pullout processes at different stages for: (a) slow cooled (1⬚C/min); and (b) fast cooled (1800⬚C/min) specimens.
functionally similar to the variations of matrix mechanical properties, suggesting strong correlations between these properties that were controlled by cooling rate and crystallinity. It is also interesting to note that the maximum embedded fibre length above which the fibre fractured without being pulled out was significantly larger (⬃90 mm) for the fast cooled specimen than for the slow cooled counterpart (⬃60 mm). It was also highlighted that the absolute values of IFSS
were consistently 20–30% higher than ILSS at a given cooling rate (Fig. 16(a)). This is not surprising in view of the differences in nature and loading conditions of the two tests. It is well known that the ILSS is sensitive to the bulk matrix properties, while the IFSS measured from fibre pullout tests is more specific to the fibre–matrix interface bond [47]. However, when these interphase properties were normalised with the respective highest values corresponding to the lowest cooling rate used in this study, i.e. 1⬚C/min, good
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Fig. 14. Optical microphotographs of a fibre pullout specimen cooled at 1800⬚C/min, showing significant plastic deformation in the matrix surrounding the fibre socket.
agreement existed between the IFSS and ILSS in terms of both the trend and magnitude with respect to cooling rate (Fig. 16(b)). The similar level of sensitivity of these properties on cooling rate further confirms that cooling rate had similar influences on the interphase adhesion and the
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mechanical properties of bulk matrix material. Obviously, the matrix crystallinity showed a much higher sensitivity than the IFSS and ILSS. This may mean there were other important parameters apart from crystallinity itself, such as lamella morphology and its role in chemical bonding, that were partly responsible for the variations of these properties. The high IFSS values at low cooling rates can be explained in terms of the molecular structure and crystalline morphology of the interphase. The dependence of fibre– matrix interface bonding mechanisms on the degree of crystallinity and the lamella is schematically shown in Fig. 17. The PEEK matrix consolidated for a long time at high temperatures tends to uncoil more completely, diffuse and adsorb more strongly onto the fibre surface than those processed at high cooling rate [48]. The PEEK chains with polar carbony1 groups start losing conformations near the fibre surface, tending to flatten out and occupy a lower energy state. The flattened crystalline lamella chains with high crystallinity at the interphase region would allow more polar components to interact with the functionalities present on the fibre surface, thereby increasing the work of adhesion and the interfacial shear strength, as suggested by Felix et al. [49] for a cellulose fibre/polypropylene system. The thickness of the oriented lamella layer adjacent to the fibre surface for a similar carbon/PEEK system was estimated to be a maximum of 5 mm under the favourable thermal condition [34]. The transcrystalline phase introduced by interface shear and temperature gradient may, to a certain extent, contributed to the high IFSS [50], although transcrystallites as opposed to spherullites were not specifically evidenced in this study. In sharp contrast, at a high cooling rate, the PEEK matrix is rather difficult to crystallise in the low super-cooling state, resulting in an amorphous PEEK-rich interphase, which is unable to develop strong bonds with the fibre surface.
Fig. 15. Maximum debonding force versus embedded fibre length for different cooling rates.
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Fig. 16. Comparisons of: (a) IFSS, ILSS, degree of crystallinity; and (b) the normalised values as a function of cooling rate.
4. Concluding remarks The fibre–matrix interphase adhesion as affected by different cooling rates was characterised for a carbon fibre/PEEK composite. The degree of crystallinity and mechanical properties of bulk matrix material, as well as the lamella morphology at the interphase region surrounding the fibre played the major role in determining the interface bond strength. Major findings from the experimental studies are highlighted in the following. There was a strong correlation between cooling rate and the crystallinity for both the neat PEEK resin and the carbon fibre/PEEK composite. The degree of crystallinity and the spherullite size decreased with increasing cooling rate, as evidenced from the DSC measurements and cross-polarised
microscopy. The low crystallinty and small spherullites at a high cooling rate were explained by the low degree of molecular perfection and the low mobility of polymer chains, which limit the ability of chains to diffuse into the growing crystal front. The crystalline structure developed at a low cooling rate was better defined and had more distinct and thicker lamellae than that developed at a high cooling rate. The tensile strength and elastic modulus of neat PEEK resin decreased, while the ductility increased with increasing cooling rate through its effect on crystallinity. The high degree of crystallinity with ordered thick lamellae crystallites obtained at a low cooling rate caused the deformation slips in the crystal blocks to be more difficult. The improvement of crystalline perfection and lamella thickening taking
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Acknowledgements This project was financially supported by the Research Grant Council of Hong Kong (CERG HKUST719/95E) and the Hong Kong University of Science & Technology (HKUST) Area of Excellence Grant (AoE97/98.EG14). Experiments were conducted with the technical supports from Advanced Engineering Materials Facility and Materials Characterisation & Preparation Facilities of HKUST. Part of the paper was presented at Third International Conference on Fracture and Strength of Solids, Hong Kong, 1997, and First Asian–Australasian Conference on Composite Materials (ACCM-1), Osaka, Japan, 1998.
References Fig. 17. Scanning electron micrographs of fracture surface and schematic fibre–matrix interaction of carbon fibre/PEEK composites cooled at: (a) 1⬚C/min; and (b) 1800⬚C/min.
place in a slow cooling process promoted the post-yield hardening behaviour as reflected by the absence of load drop after plastic yielding. The interphase bond strength decreased with increasing cooling rate through its effects on crystallinty and lamella morphology of interphase. The crystallinity in the matrix controlled the mechanical properties of bulk matrix, including the ductility, tensile strength and elastic modulus, while the lamella morphology on the fibre surface affected the interface adhesion mechanisms. The interface failure mechanisms were also a direct reflection of the interface adhesion mechanisms. The closely packed crystalline lamella interphase that was adsorbed more strongly onto the fibre surface in slow cooled specimens gave rise to a high IFSS with relatively brittle interface debonding. In contrast, the amorphous PEEK-rich interphase introduced in fast cooled specimens resulted in a low IFSS and a relatively ductile interface failure, as manifested by the thick permanently deformed matrix layer surrounding the pulled out fibre socket. The interphase adhesion represented by IFSS and ILSS, and the bulk mechanical properties of matrix, such as ductility, strength and modulus, all exhibited functionally similar variations with cooling rate. The sharp transition of these properties at about 600⬚C/min may be attributed to a transition of failure mechanisms associated with the degree of crystallinity. The sensitivity of matrix crystallinity on cooling rate was much higher than the changes in IFSS and ILSS. In summary, it is demonstrated in this study that the fibre–matrix interphase adhesion and the bulk mechanical properties of the thermoplastic composite can be optimised, if not totally maximised, by controlling the crystallinity which is affected by processing conditions, such as the cooling rate.
[1] Lustiger A, Uralil FS, Newaz GM. Processing and structural optimization of PEEK composites. Polym Compos 1990;11:65–75. [2] Saiello S, Kenny J, Nicolais L. Interface morphology of carbon fibre/ PEEK composites. J Mater Sci 1990;25:3493–6. [3] Zhang M, Xu J, Zhang Z, Zeng H, Xiong X. Effect of transcrystallinity on tensile behaviour of discontinuous carbon fibre reinforced semicrystalline thermoplastic composites. Polymer 1996;37:5151–8. [4] Kim JK, Mai YW. Interfaces in composites. In: Chou TW, editor. Materials science and technology, structure and properties in composites, 13. Weinheim, Germany: VCH, 1993. pp. 239–89. [5] Youssef Y, Denault J. Thermoformed glass fiber reinforced polypropylene: microstructure, mechanical properties and residual stresses. Polym Compos 1998;19:301–9. [6] Waterbury MC, Drzal LT. Interfacial shear strengths of carbon fibres in bisphenol-A polycarbonate. In: Ishida H, editor. Controlled interphases in composite materials, New York: Elsevier, 1990. pp. 731–9. [7] Nielsen AS, Pyrz R. Study of the influence of thermal history on the load transfer efficiency and fibre failure in carbon/polypropylene microcomposites using Raman spetroscopy. In: Proceedings of Seventh Int Conf Compos Interfaces (ICCI-VII), Shonan, Japan, 1998, paper SXI-2. [8] Schulz E, Kalinka G, Auersch W. Effect of transcrystallization in carbon fiber reinforced polyphenylene sulfide composites on the interfacial shear strength investigated with the single fiber pull-out test. J Macromol Sci: Phys B 1996;35:527–46. [9] Ye L, Scheuring T, Friedrich K. Matrix morphology and fibre pull-out strength of T700/PPS and T700/PET thermoplastic composites. J Mater Sci 1995;30:4761–9. [10] Peron B, Lowe A, Baillie C. The effect of transcrystallinity on the interfacial characteristics of polypropylene/alumina single fibre composites. Composites Part A 1996;27:839–45. [11] Marom G. The microstructure of crystalline interfaces and effects on properties of fibre-reinforced composites. In: Proc Seventh Intern Conf Compos Interfaces (ICCI-VII), Shonan, Japan, 1998, paper SVII-1. [12] Beehag A, Ye L. Fibre/matrix adhesion in thermoplastic composite: is transcrystallinity a key? In: Proc Eleventh Int Conf Compos Mater (ICCM-11), Gold Coast, Australia, 1997;4:723–730. [13] Zou YL, Netravali AN. Ethylene/ammonia plasma polymer deposition for controlled adhesion of graphite fibers to PEEK. II. Effect on fiber and fiber/matrix interface. J Adhes Sci Technol 1995;9:1505– 20. [14] Kobayashi H, Hayakawa E, Kikutani T, Takaku A. Effect of quenching and annealing on fiber pull-out from crystalline polymer matrices. Adv Compos Mater 1991;1:155–68. [15] Folkes MJ, Kalay G, Ankara A. The effect of heat treatment on the properties of PEEK and APC2. Compos Sci Technol 1993;46:77–83.
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S.-L. Gao, J.-K. Kim / Composites: Part A 31 (2000) 517–530
[16] Gao SL, Kim JK. Interphase morphology and fibre pull-out behaviour of carbon fibre-PEEK composites. Key Engng Mater 1998;145149:811–6. [17] Gao SL, Kim JK. Effect of cooling rate on impact damage resistance of carbon fibre-PEEK laminates. In: Proc Eighth European Conf Compos Mater (ECCM-8), Naples, Italy, 1998:95–104. [18] Gao SL, Kim JK. Effect of cooling rate on interphase properties and mechanical response of carbon fibre/PEEK composites. Mater Sci Res Int 1999;5:157–62. [19] Blundell DJ, Osborn BN. The morphology of poly(aryl-ether-etherketone). Polymer 1983;24:953–8. [20] Kim JK, Zhou LM, Mai YW. Techniques for studying composite interfaces. In: Cheremisinoff NP, editor. Handbook of advanced materials testing, New York: Marcel Dekker, 1995. pp. 327–65. [21] Ohsawa T, Nakyama A, Miwa M, Hasegawa A. Temperature dependence of the critical fibre length for the glass fibre reinforced thermosetting resins. J Appl Polym Sci 1978;22:3203–17. [22] Drzal LT. Fibre–matrix interphase structure and its effect on adhesion and composite mechanical properties. In: Ishida H, editor. Controlled interphases in composite materials, New York: Elsevier, 1990. pp. 309–20. [23] Vu-Khanh T, Denault J. Effect of molding parameters on the interfacial strength in PEEK/carbon composites. J Reinf Plast Compos 1993;12:916–31. [24] Blundell DJ, Chalmers JM, Mackenzie MW, Gaskin WF. Crystalline morphology of the matrix of PEEK-carbon fibre aromatic polymer composites. I. Assessment of crystallinity. SAMPE Quart 1985;16:22–30. [25] Chen M, Chung CT. Crystallinity of isothermally and nonisothermally crystallized poly(ether ether ketone) composites. Polym Compos 1998;19:689–97. [26] Xiao XR, Hoa SV. Effect of melting history on the crystalline characteristics of polyetheretherketone aromatic polymer composite. Theor Appl Fract Mech 1990;14:49–56. [27] Davies P, Cantwell WJ, Jar PY, Richard H, Neville DJ, Kausch HH. Cooling rate effects in carbon fibre/PEEK composites. In: O’Brien TK, editor. Composite materials: fatigue and fracture, ASTM STP 1110, Philadelphia, 1991. pp. 70–88. [28] Talbott MF, Springer GS, Berglund LA. The effects of crystallinity on the mechanical properties of PEEK polymer and graphite fiber reinforced PEEK. J Compos Mater 1987;21:1056–81. [29] Velisaris CN, Seferis JC. Crystallization kinetics of polyetheretherketone (PEEK) matrices. Polym Engng Sci 1986;26:1574–81. [30] Unger WJ, Hansen JS. The effect of thermal processing on residual strain development in unidirectional graphite fibre reinforced PEEK. J Compos Mater 1993;27:59–82. [31] Yasuniwa M, Nakafuku C. Polymer 1980;12:105. [32] Partridge IK, Davies P, Parker DS, Yee AF. Yield and fracture in PES and PEEK matrix polymers and their composites. In: Proc Int Conf Polym Compos, Cranfield, UK: Plastics and Rubber Inst, 1987;5:1– 12.
[33] Uralil FS, Newaz GM, Lustiger A. Processing effects and damage tolerance in poly(etheretherketone) composites. Polym Compos 1992;13:7–14. [34] Lee Y, Porter RS. Crystallization of polyetheretherketone (PEEK) in carbon fibre composites. Polym Engng Sci 1986;26:633–9. [35] Waddon AJ, Hill MJ, Keller A, Blundell DJ. On the crystal texture of linear polyaryls (PEEK, PEK and PPS). J Mater Sci 1987;22:1773– 84. [36] Chen EJH, Hsiao BS. The effects of transcrystalline interphase in advanced polymer composites. Polym Engng Sci 1992;32:280–6. [37] Deporter J, Baird DG. The effects of thermal history on the structure/ property relationship in polyphenesulfide/carbon fibre composites. Polym Compos 1993;14:201–13. [38] Tregub A, Harel H, Marom G, Migliaresi C. The influence of thermal history on the mechanical properties of poly(ether ether ketone) matrix composite materials. Compos Sci Technol 1993;48:185–90. [39] Jar PY, Kausch HH. Annealing effect on mechanical behaviour of PEEK. J Polym Sci Polym Phys 1992;30:775–8. [40] Weibull W. A statistical distribution function of wide applicability. J Appl Mech 1951;18:293–6. [41] Galiotis C, Young RJ, Yeung PHJ, Batchelder DN. The study of model polydiacetylene/epoxy composites. I. The axial strain in the fibre. J Mater Sci 1984;19:3640–8. [42] Asloun EM, Nardin M, Schultz J. Stress transfer in single-fibre composites: effect of adhesion, elastic modulus of fibre and matrix, and polymer chain mobility. J Mater Sci 1989;24:1835–44. [43] Ogata N, Yasumoto H, Yamasaki K, Hu Y, Ogihara T, Yanagawa T, Yoshida K, Yamada Y. Evaluation of interfacial properties between carbon fibres and semicrystalline thermoplastic matrices in singlefibre composites. J Mater Sci 1992;27:5108–12. [44] Cox HL. The elasticity and strength of paper and other fibreous materials. Brit J Appl Phys 1952;3:72–79. [45] Drzal LT, Rich MJ, Koenig MF, Lloyd PF. Adhesion of graphite fibres to epoxy matrices: II. Effect of fibre finish. J Adhes 1983;16:133–52. [46] Pisanova EV, Zhandarov SF. On the mechanism of failure in microcomposites consisting of single glass fibres in a thermoplastic matrix. Compos Sci Technol 1997;57:937–43. [47] Kim JK, Mai YW. Engineered interfaces in fibre reinforced composites. Oxford, UK: Elsevier, 1998 p. 43–92. [48] Brady RL, Porter RS. Interfacial adsorption and crystallization of polycarbonate in carbon fibre composites. J Appl Polym Sci 1990;39:1873–85. [49] Felix JM, Gatenholm P. Effect of transcrystalline morphology on interfacial adhesion in cellulose/polypropylene composites. J Mater Sci 1994;29:3043–9. [50] Clark RLJ, Kander RG, Sauer BB. Nylon 66/poly(vinyl pyrrolidone) reinforced composites. I. Interphase microstructure and evaluation of fiber–matrix adhesion. Composites Part A 1999;30:27–36.
Cooling rate influences in carbon fibre/PEEK composites. Part 1. Crystallinity and interface adhesion Shang-Lin Gao, Jang-Kyo Kim* Department of Mechanical Engineering, Hong Kong University of Science and Technology, Clear Water Bay, Hong Kong, People’s Republic of China Received 17 June 1999; received in revised form 22 November 1999; accepted 19 January 2000
Abstract The effect of cooling rate on the fibre–matrix interface adhesion for a carbon fibre/semicrystalline polyetheretherketone (PEEK) composite was characterised based on the fibre fragmentation, fibre pullout and short beam shear tests. The interface adhesion was correlated to the degree of crystallinity and the crystalline morphology, as well as the bulk mechanical properties of neat PEEK resin, all of which were in turn controlled by cooling rate. It was shown that the interface bond strength decreased with increasing cooling rate; the tensile strength and elastic modulus of PEEK resin decreased, while the ductility increased with increasing cooling rate through its dominant effect on crystallinity and spherullite size. The improvement of crystalline perfection and flattened lamella chains with high crystallinity at the interphase region were mainly responsible for the strong interface bond in composites processed at a low cooling rate. The interphase failure was characterised by brittle debonding in slow-cooled composites, whereas the amorphous PEEK-rich interphase introduced in fast cooled specimens failed in a ductile manner with extensive plastic yielding. 䉷 2000 Elsevier Science Ltd. All rights reserved. Keywords: Cooling rate; Strength; B. Mechanical properties
1. Introduction Thermoplastic matrix composites have been studied extensively in an effort to best utilise the potential advantages of high fracture toughness, high temperature resistance, repairability and ease of manufacture. It has long been recognised that the mechanical properties of semicrystalline polymers and the composites made therefrom are dependent on the crystallinity and crystalline morphology, which are strongly affected by the processing conditions [1,2]. The presence of carbon fibres within the matrix induces nucleation and growth of crystallites perpendicular to the fibre surface, i.e. transcrystallisation, which has a considerable influence on the fibre/matrix interface interaction and the failure behaviour in both the matrix and the interphase region [3]. Optimisation of the bulk mechanical properties of thermoplastic matrix composites, therefore, requires better understanding of the effect of processing conditions on the interphase properties. The interactions at the interphase region in semicrystalline thermoplastic composites depend on a number of
* Corresponding author. Tel.: ⫹ 852-2358-7207; fax: ⫹ 852-23581543. E-mail address: [email protected] (J.-K. Kim).
factors, such as matrix morphology, fibre surface condition, presence of residual stresses, moduli of the fibre and matrix, as well as the presence of reactive functionalities [4]. The majority of these characteristics are determined by processing conditions, including moulding temperature, cooling rate, holding time/temperature and annealing conditions. A great deal of studies have been directed toward understanding the influences of these parameters, but the results reported hitherto displayed no apparent consensus. For example on the effect of cooling rate, an increase in the interfacial shear bond strength (IFSS) was noted with increasing cooling rate for glass fibre/polypropylene (PP) [5] and carbon fibre/polycarbonate (PC) systems [6]. This result was attributed to the improved wetting behaviour, high ductility, fracture energy and low residual stresses in the matrix and the inhibition of molecular weight segregation or weak boundary layer that arose from the high cooling rate. However, there were contradictory reports for carbon fibre/PP matrix [7], carbon fibre/polyphenylene sulfide (PPS) [8], carbon fibre/polyethylene terephthalate (PET) [9] and alumina fibre/PP systems [10], showing a totally opposite trend, or no significant cooling rate dependence. The high IFSS with decreasing cooling rate was partly due to the formation of transcrystalline interface and partly to the improved adsorption capability.
1359-835X/00/$ - see front matter 䉷 2000 Elsevier Science Ltd. All rights reserved. PII: S1359-835 X( 00)00 009-9
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Table 1 Reported interface shear strength of carbon fibre/PEEK composites System
IFSS (MPa)
Testing methods
Processing condition
Reference
AS4/PEEK AS4/PEEK AS4/PEEK AS4/PEEK
38–43 50–97 70–100 “Good interface adhesion for slow cooling”
Pullout Fragmentation Pullout SEM observation
Isothermal crystallisation Cooling rate: ⬃400⬚C/min Quenching and annealing Cooling rate: 2–63⬚C/min
[9] [13] [14] [15]
The effects of transcrystallinity on the IFSS and the bulk mechanical properties of composite were also found to be inconsistent. It was argued that large thermal stresses might build up during the transcrystallisation process, which in turn reduced the IFSS, whereas the increasing degree of crystallinity of matrix material effectively decreased the mismatch in moduli between the fibre and matrix, and thus enhancing the stress transfer efficiency across the interface. The formation of such transcrystalline layer, thus, improved the tensile modulus, strength and fracture toughness of short carbon fibre/PEEK matrix composites [3]. In contrast, Marom [11] attributed the effect of transcrystalline layer to a preferred crystallite orientation relative to the fibre, thereby providing the surrounding matrix with a high rigidity and reduced thermal expansion in the fibre direction, which effectively lowered the residual thermal stresses. Meanwhile, transcrystallinity was found to have conflicting influences on the IFSS, depending on the matrix material [12]. For carbon fibre/PEEK systems, the presence of transcrystalline interphase layer gave rise to the IFSS, whereas for carbon/PES system, the reverse was true, although the differences were only marginal in both cases. Even for an identical AS4 carbon fibre/PEEK system with which the present study is mainly concerned, the measured IFSS varied widely from 40 to 110 MPa, due mainly to the largely different processing conditions [12–15], as presented in Table 1. The use of different testing techniques and different data reduction methods were also partly responsible for the large data variation. The inconsistency on the effects of cooling rate and crystallinity warrants a systematic study to clarify the relationship between cooling rate, morphology and the IFSS. The present study is the continuation of our previous work [16–18] on the improvement of mechanical performance and damage tolerance of carbon fibre/ PEEK matrix composites through the optimisation of processing conditions and control of the fibre–matrix interphase properties. The dependence of IFSS on matrix crystallinity was evaluated over a wide range of cooling rate from 1 to 2000⬚C/min, using a variety of thermomechanical characterisation techniques and mechanical testing methods. Model microcomposites were employed to measure the IFSS in the single fibre pullout and fragmentation tests.
2. Experimental 2.1. Materials and characterisation of crystallinity All specimens used in the present study were prepared from Hercules AS4 carbon fibres, PEEK powder (Victrex) and APC-2 prepregs containing continuous AS4 carbon fibre of 61% by volume. The PEEK powder was a similar material to the matrix of APC-2 prepregs according to the manufacturer (ICI Fiberite). The crystallinity of both neat PEEK resin and carbon fibre/PEEK composites was measured using a differential scanning calorimeter (DSC, SETAPRAM-92). Specimens of approximately 10 mg were analysed in nitrogen at a heating rate of 20⬚C/min. The following equation was used to calculate the degree of crystallinity, X: X
DHm ⫺ DHc DHf 1 ⫺ a
1
where DHm and DHc are the enthalpy of fusion at melting point as measured by the area under the endothermic peak, and the enthalpy of crystallisation as measured by the area under the exothermic crystallisation peak, respectively. DHf is the enthalpy of fusion of fully crystalline PEEK, which is taken 130 J/g [19]; and a is the mass fraction of carbon fibre in the composite. An optical microscope equipped with polarised light filters was used to characterise the crystalline morphology of the interphase and the surrounding matrix. The surface of metallographically polished composite specimens was etched using a mixture of 1 wt% potassium permanganate, surfuric acid (five parts), ortho-phosphoric acid (two parts) and distilled water (two parts by volume). 2.2. Single fibre pullout and fragmentation tests The fibre pullout and fibre fragmentation tests are amongst the most popular techniques to measure the fibre–matrix interface bond quality [20]. Along with the usual difficulties in measuring the interfacial shear strength based on single fibre model composites because of the small diameter of carbon fibres, the PEEK matrix, in particular, presented problems in specimen fabrication due to the high melt viscosity and high melting temperature. These problems necessitated the development of special methods to fabricate single fibre composite specimens at different
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different thickness depending on the desired cooling rate, which were then cooled between two aluminium plates. Cooling rates above 1000⬚C/min were achieved by directly quenching the moulds in a water bath of ambient temperature. The moulds were wrapped using aluminium foil to avoid direct contact with water. The cooling rate was monitored using a Datapag tracker with thermocouples embedded in the specimen. The pullout and fragmentation experiments were performed on a Minimat testing machine at a cross-head speed of 0.01 mm/min and 10 mm/min, respectively. The fibre/matrix interface morphology and the fibre fragment lengths were simultaneously monitored using a polarised light microscope, a video recorder and an image analyser. The whole testing set-up is schematically shown in Fig. 2. The critical transfer length, lc, was calculated from the measured mean fibre fragment lengths, l [21]: lc 4l=3
Fig. 1. Schematic of specimen preparation for: (a) single-fibre pullout; and (b) fibre fragmentation tests.
cooling rates. Specimens were prepared by partially or wholly embedding a carbon fibre within the PEEK powder, which was contained in a mould, as shown schematically in Fig. 1. The moulds were heated to 400⬚C in an oven for 10 min, which were then allowed to cool as detailed in the following. Cooling rates below 10⬚C/min were achieved by cooling the moulds in a preheated oven of different initial temperatures. Moulds for cooling rates between 10 and 1000⬚C/min were shielded using polyimide films of
2
The critical aspect ratio, lc/d and the modulus ratio, Ef/Em, of fibre to PEEK matrix were calculated for the AS4 carbon fibre of an average diameter d 7:7 mm and modulus Ef 234:4 GPa [22]. In view of the fact that the maximum loads, F, measured from the fibre pullout tests consistently displayed a linear dependence on the embedded fibre length, L, the IFSS was calculated using the simple linear relation: IFSS
F pdL
3
The baseline mechanical properties of neat PEEK resin were measured in uniaxial tension of PEEK films according to the specification ASTM D11708. The interlaminar shear strengths were also measured from the short beam shear (SBS) tests of unidirectional carbon fibre/PEEK
Fig. 2. Schematic of Minimat testing fixture placed on a microscope stage.
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Fig. 3. DSC thermograms for: (a) neat PEEK resin; and (b) carbon fibre/PEEK composites processed at different cooling rates.
composite laminates that were fabricated from APC-2 prepregs and processed at different cooling rates. The SBS tests were conducted on a MTS 858 universal testing machine at a cross-head speed of 1.3 mm/min, according to the specification ASTM D2344. The fracture surfaces of the composites were examined using a scanning electron microscope (Cambridge Stereoscan 200).
3. Results and discussion 3.1. Crystallinity and morphology The crystallinity contents were measured using DSC and the corresponding DSC thermograms are shown in Fig. 3 for the neat PEEK resin and carbon fibre/PEEK composites
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Table 2 Effect of cooling rate on crystallinity of neat PEEK resin and carbon fibre/PEEK composite Cooling rate (⬚C/min) 1 70 160 600 1000 1500 a b
Tc (⬚C) – – – 188 (182) 187 188 (183)
DHc (J/g)
Tm (⬚C) 351 (346) 351 (344) 352 352 (342) 352 351 (342)
a
– – – 3 具3典 13 14 具15典
DHm (J/g) 49 38 37 37 38 36
具42典 具36典 具36典 具31典
b
Crystallinity (%) 38 具33典 30 具28典 28 26 具25典 19 17 具12典
Values in parentheses ( ) represent those for carbon fibre/PEEK composite. Values in parentheses 具 典 represent those for carbon fibre/PEEK composite taking into account the matrix weight fraction of 32%.
processed at different cooling rates. The PEEK resin and the composite processed at a very high cooling rate, say at 1000⬚C/min or above, showed only a marginal difference between the areas of crystallinity peak and melting peak, indicating a very low degree of crystallinity. The semicrystalline PEEK resin and the composite processed at a low cooling rate displayed almost a negligible crystallisation peak, giving rise to a high degree of crystallinity. The results are summarised in Table 2 and the crystallinity is plotted as a function of cooling rate in Fig. 4. Also included in this figure is the compilation of crystallinity values reported in the literature [23–29]. The data indicates considerable scattering, in part attributable to the different techniques employed to measure crystallinity, such as DSC, wideangle X-ray spectroscopy (WAXS) and the density method [30], as well as to the different types of fibres and matrix materials studied. Even so, the general trend clearly indicates that the higher the cooling rate, the lower the crystallinity. For both the neat PEEK resin and carbon fibre/PEEK composites, the degree of crystallinity decreased as the cooling rate increased (Fig. 4). It is well known that
crystallinity, melting point and cooling rate are closely interrelated. The melting peak temperature of a semicrystalline polymer composite is determined by several factors, such as crystal size, fold surface energy and plate thickness [31]. The size of sphrullite in turn is strongly influenced by cooling rate: the lower the cooling rate, the larger the spherullite size with a higher degree of molecular perfection and a lower melting temperature [27,32,33]. The correlation between the spherullite size and cooling rate is clearly seen from the cross-polarised optical micrographs given in Fig. 5. The spherullites in the PEEK resin obtained after very slow cooling, say 1⬚C/min, was as large as a few hundred mm in diameter, while they were too small to be detected using an optical microscope once a very high cooling rate was used, say 2000⬚C/min. The crystallite size dependence on cooling rate can be explained from the viewpoint of molecular movement: the mobility of polymer chains is discouraged at a high cooling rate, thus limiting the ability of the chains to diffuse into the growing crystal front. Another significant implication of Fig. 5 is that the fibre surface influenced the nucleation and growth of spherullites within the matrix, particularly at a low cooling rate. It is
Fig. 4. Degree of crystallinity as a function of cooling rate for PEEK resin and carbon fibre/PEEK composites. Measurement methods are shown in ( ).
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Fig. 6. (a) Scanning electron micrograph of fibre-nucleated spehrullites on an etched cross-section of a slowly cooled carbon fibre/PEEK composite. (b) Schematic diagram showing the growth of highly oriented spherullites from the fibre surface [36]. Fig. 5. Optical micrographs of PEEK matrix morphologies around isolated carbon fibres at cooling rates: (a) 1⬚C/min; (b) 200⬚C/min; (c) 1000⬚C/min; and (d) ⬃2000⬚C/min.
well known that nucleation and growth of spherullites from the surface of high modulus carbon fibres is often related to the graphitic nature of the fibre surface rather than the fibre surface chemistry [34,35]. A further proof for these fibre surface-induced spherullites is presented in Fig. 6. The spherullites emerging from the fibre surface impinged each other within the narrow matrix bed due to the presence of rigid fibres of high volume fraction. The schematic illustrated in Fig. 6(b) [36] depicts spherullites of a high order of alignment and compact packing. It is believed that the crystalline structure developed at a slow cooling rate is better defined and larger in size with more distinct and thicker lamellae than that developed at a high cooling rate. The fast-cooled composite generally displayed a mixture of isolated and fibre-nucleated spehrullites, whereas
Fig. 7. Tensile stress–strain curves of neat PEEK resin processed at different cooling rates.
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Fig. 8. Ductility, tensile strength and modulus of neat PEEK as a function of cooling rate.
the slow-cooled composite only revealed much larger fibre nucleated spherullites [1]. It is interesting to note that both the degree of crystallinity and the melting peak temperature of the composite were slightly lower than those of the neat PEEK resin for the same processing condition (Table 2). A similarly reduction in melting temperature for a carbon/PPS composite [37] and in crystallinity for a carbon/PEEK composite [38] were previously reported. It seems that the presence of densely packed fibres within the matrix severely suppressed the growth of spherullites, effectively offsetting the aforementioned ameliorating effect of fibre surface acting as preferred crystallite nucleation sites. 3.2. Tensile properties of matrix Typical engineering tensile stress-strain curves of neat PEEK resin are presented in Fig. 7, and the average values of the major mechanical properties are summarised in Fig. 8. It is obvious that both the elastic modulus and the strength of PEEK increased, while the ductility decreased with
523
decreasing cooling rate. As would be expected from the DSC results (Fig. 4), the high crystallinity with ordered crystallites obtained at a low cooling rate gave rise to a high stiffness of the material, causing yielding to be more difficult with dramatically reduced ductility. A high stiffness and high strength means that it is more difficult to induce slips in the crystal blocks due to the higher orderliness and thicker lamellae. Fig. 7 further indicates that there were distinct stress drops after yield point in the specimens cooled at a high cooling rate, say at or above 1000⬚C/min, while the stress drop was totally absent for the specimens processed at low cooling rates. Similar tensile behaviour for the annealed PEEK resin suggests [39] that the stress drop at the yield point was hidden by the occurrence of postyield hardening behaviour due mainly to the improvement of crystalline perfection and lamella thickening taking place during the annealing treatment. It is expected that a slow cooling process promotes the formation of more perfect crystallites with a higher proportion of thick lamellae than a fast cooling process. It is also worth noting that an increase of cooling rate from 1⬚C/min to about 600⬚C/min resulted in a significant change in all properties measured, including the ductility, tensile strength and modulus. Further increase in cooling rate beyond 600⬚C/min did not much vary these values. This distinct behaviour may be a direct reflection of the change in crystallinity, and similar trends with respect to cooling rate were noted for the IFSS and the composite interlaminar shear strength (ILSS), as discussed in Section 3.3.2. 3.3. Interaction at the interphase region 3.3.1. Fibre fragmentation test There was considerable scattering in the fibre fragment length, l. Thus, the statistical variation of l was described by the cumulative probability, P, which was calculated from: P l
n l N ⫹ 1
4
where n(l) is the number of fibres whose lengths are equal to or shorter than l, and N is the total number of fragmented fibres at saturation. The experimental data were analysed using the simple cumulative probability theory proposed previously [40]: P l 1 ⫺ exp⫺ l=l0 m
Fig. 9. Cumulative fibre fragment length for different cooling rates.
5
where m and l0 are the Weibull modulus (or shape factor) and the characteristic length (or scale factor) of fragmented fibres, respectively. A strong correlation was noted between the fibre fragment length distribution and cooling rate. The cumulative distributions are presented in Fig. 9, and the Weibull parameters as well as the mean fibre fragment lengths are summarised in Table 3. It is obvious from these results that a shift to the shorter fragment lengths corresponded to a slower cooling rate. Meanwhile, the
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Table 3 Weibull parameters for fibre fragmentation lengths at saturation Cooling rate (⬚C/min)
Weibull modulus, m
Weibull scale parameter lo (mm)
Mean fibre fragmentation length ^ SD (mm)
490 600 1000 1500 1800
2.88 2.81 3.55 3.16 2.91
399 487 525 550 610
350 ^ 45 432 ^ 82 470 ^ 110 491 ^ 64 543 ^ 114
Weibull modulus, m, varied rather irregularly with the cooling rate without showing any general trend, indicating that the distribution of fibre fragment lengths was insensitive to cooling rate. The critical fibre aspect ratio, lc/d, calculated from the mean fragment length at saturation is plotted as a function of cooling rate in Fig. 10. There is approximately a linear relation between these parameters: the higher the cooling rate, the higher the critical aspect ratio. This indicates that the efficiency of stress transfer across the fibre–matrix interface is inversely proportional to cooling rate. The high stress transfer efficiency at a low cooling rate may be explained in terms of improved mechanical characteristics of the matrix material and, to a certain extent, the high interface bond quality. In fact, the major factors behind the variation of stress transfer efficiency measured by fibre fragment lengths are rather matrix-related: the elastic modulus of the matrix material played a dominant role. The critical aspect ratio, lc/d, is plotted as a function of the modulus ratio of fibre to matrix, Ef/Em, in Fig. 11, clearly indicating there is a strong correlation between these parameters. This again confirms that the degree of crystallinity and crystalline morphology were mainly responsible for the variation of stress transfer efficiency through the effect of cooling rate. The above observation is quite consistent with experimental evidence on other composite systems that lc/d varied approximately
with (Ef/Em) 1/2 [41–43], as the early shear lag model suggests [44]. It should be mentioned here that Figs. 9 and 10 contain only the data corresponding to the specimens that were processed at high cooling rates and were loaded until saturation of fibre fragmentation. The specimens processed at cooling rates lower than about 400⬚C/min, however, tended to fail prematurely due to the cracks originating from the broken fibre ends, disallowing the measurement of saturation fibre fragment lengths. Also associated with this observation is the cooling rate-dependent failure behaviour of the fibre–matrix interface near the broken fibre ends. Fig. 12 presents typical microphotographs showing distinct failure modes. A transverse matrix crack in a peculiar diamondshape is clearly seen for the specimen with a stiff matrix and/or a strong interface, corresponding to a low cooling rate. Similar transverse cracks appeared at almost all broken fibre ends. In contrast, a gap between two broken fibres is detected for the specimens with a compliant, ductile matrix and/or a weak interface, corresponding to a high cooling rate. With increasing cooling rate, the failure mode changed from matrix crack propagation to interfacial debonding, with the empty cylindrical channel representing a debonded interface [45,46]. The failure mode change is a direct reflection of the improvement in ductility and fracture resistance of the matrix.
Fig. 10. Critical fibre aspect ratio, lc/d, versus cooling rate for carbon fibre/PEEK composites.
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Fig. 11. Critical fibre aspect ratio, lc/d, versus modulus ratio of fibre to matrix, Ef/Em for carbon fibre/PEEK composites.
3.3.2. Fibre pullout test The results obtained from the fibre pullout test complement and, most often, support the implications derived from the fibre fragmentation test. The interface failure mechanisms taking place during the fibre pullout processes were studied based on the sequential polarised microphotographs and the typical pullout load–displacement records, as shown in Fig. 13. The interface debonding process in the specimen processed at a very low cooling rate, e.g. 1⬚C/min (Fig. 13(a)), was characterised by brittle fracture with little matrix deformation along the fibre. As a consequence, almost a linear increase in pullout load was noted until it reached the maximum. In the fast cooled specimens at 1800⬚C/min (Fig. 13(b)), however, the rising part of the load–displacement trace exhibited an abrupt change in slope, indicating a rather stable interphase failure process between the two peaks. A typical failure mechanism was ductile yielding of the interphase, whose details are discussed in the
Fig. 12. Photomicrographs of stretched carbon fibre/PEEK fragmentation specimens cooled at: (a) 500⬚C/min; and (b) 1800⬚C/min.
following. At stages 1–3, the interphase deformed viscoelastically until yielding took place, as indicated by the bright fringe patterns around the fibre surface. At stages 3–5, interface debonding propagated stably from the embedded fibre end along with extensive plastic yielding of the interphase. At stages 5–6, the interface crack propagated unstably, and the fibre was completely pulled out. A thick layer of permanently deformed material is clearly seen surrounding the fibre socket after the complete pullout (see the photograph for stage 6). It appears that yielding took place at the early stage of pullout, and the thickness of the matrix layer affected by plastic deformation was approximately twice the fibre diameter for a typical specimen, as shown in Fig. 14. Examination of the fracture surfaces of pulled out fibres revealed that: for the slow cooled specimens, the fibre was relatively clean and free of matrix residue, indicating rather brittle interfacial debonding. In contrast, for the fast cooled specimens, some residual pieces of PEEK resin were seen adhering to the fibre surface, indicating plastic deformation of the soft, ductile matrix before interphase failure. Having studied the interphase phenomena during the fibre pullout test, which were shown to be very sensitive to cooling rate, the maximum pullout-force data were analysed in the following. The maximum pullout-force taken from the load–displacement curves were found to increase linearly with embedded fibre length for all specimens processed at different cooling rates, as shown in Fig. 15. The IFSS values calculated based on Eq. (3) are presented in Fig. 16 along with the results from the short beam shear test and DSC analysis. An increase in cooling rate from 1⬚C/min to about 600⬚C/min resulted in a significant reduction in all properties measured, including IFSS, ILSS and the degree of crystallinity. Further increase of cooling rate beyond 600⬚C/min, however, did not much vary the IFSS and ILSS, while the degree of crystallinity decreased further with cooling rate. Comparison of Figs. 8 and 16(a) revealed that the dependence of IFSS and ILSS on cooling rate was
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Fig. 13. Sequential microphotographs showing fibre pullout processes at different stages for: (a) slow cooled (1⬚C/min); and (b) fast cooled (1800⬚C/min) specimens.
functionally similar to the variations of matrix mechanical properties, suggesting strong correlations between these properties that were controlled by cooling rate and crystallinity. It is also interesting to note that the maximum embedded fibre length above which the fibre fractured without being pulled out was significantly larger (⬃90 mm) for the fast cooled specimen than for the slow cooled counterpart (⬃60 mm). It was also highlighted that the absolute values of IFSS
were consistently 20–30% higher than ILSS at a given cooling rate (Fig. 16(a)). This is not surprising in view of the differences in nature and loading conditions of the two tests. It is well known that the ILSS is sensitive to the bulk matrix properties, while the IFSS measured from fibre pullout tests is more specific to the fibre–matrix interface bond [47]. However, when these interphase properties were normalised with the respective highest values corresponding to the lowest cooling rate used in this study, i.e. 1⬚C/min, good
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Fig. 14. Optical microphotographs of a fibre pullout specimen cooled at 1800⬚C/min, showing significant plastic deformation in the matrix surrounding the fibre socket.
agreement existed between the IFSS and ILSS in terms of both the trend and magnitude with respect to cooling rate (Fig. 16(b)). The similar level of sensitivity of these properties on cooling rate further confirms that cooling rate had similar influences on the interphase adhesion and the
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mechanical properties of bulk matrix material. Obviously, the matrix crystallinity showed a much higher sensitivity than the IFSS and ILSS. This may mean there were other important parameters apart from crystallinity itself, such as lamella morphology and its role in chemical bonding, that were partly responsible for the variations of these properties. The high IFSS values at low cooling rates can be explained in terms of the molecular structure and crystalline morphology of the interphase. The dependence of fibre– matrix interface bonding mechanisms on the degree of crystallinity and the lamella is schematically shown in Fig. 17. The PEEK matrix consolidated for a long time at high temperatures tends to uncoil more completely, diffuse and adsorb more strongly onto the fibre surface than those processed at high cooling rate [48]. The PEEK chains with polar carbony1 groups start losing conformations near the fibre surface, tending to flatten out and occupy a lower energy state. The flattened crystalline lamella chains with high crystallinity at the interphase region would allow more polar components to interact with the functionalities present on the fibre surface, thereby increasing the work of adhesion and the interfacial shear strength, as suggested by Felix et al. [49] for a cellulose fibre/polypropylene system. The thickness of the oriented lamella layer adjacent to the fibre surface for a similar carbon/PEEK system was estimated to be a maximum of 5 mm under the favourable thermal condition [34]. The transcrystalline phase introduced by interface shear and temperature gradient may, to a certain extent, contributed to the high IFSS [50], although transcrystallites as opposed to spherullites were not specifically evidenced in this study. In sharp contrast, at a high cooling rate, the PEEK matrix is rather difficult to crystallise in the low super-cooling state, resulting in an amorphous PEEK-rich interphase, which is unable to develop strong bonds with the fibre surface.
Fig. 15. Maximum debonding force versus embedded fibre length for different cooling rates.
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Fig. 16. Comparisons of: (a) IFSS, ILSS, degree of crystallinity; and (b) the normalised values as a function of cooling rate.
4. Concluding remarks The fibre–matrix interphase adhesion as affected by different cooling rates was characterised for a carbon fibre/PEEK composite. The degree of crystallinity and mechanical properties of bulk matrix material, as well as the lamella morphology at the interphase region surrounding the fibre played the major role in determining the interface bond strength. Major findings from the experimental studies are highlighted in the following. There was a strong correlation between cooling rate and the crystallinity for both the neat PEEK resin and the carbon fibre/PEEK composite. The degree of crystallinity and the spherullite size decreased with increasing cooling rate, as evidenced from the DSC measurements and cross-polarised
microscopy. The low crystallinty and small spherullites at a high cooling rate were explained by the low degree of molecular perfection and the low mobility of polymer chains, which limit the ability of chains to diffuse into the growing crystal front. The crystalline structure developed at a low cooling rate was better defined and had more distinct and thicker lamellae than that developed at a high cooling rate. The tensile strength and elastic modulus of neat PEEK resin decreased, while the ductility increased with increasing cooling rate through its effect on crystallinity. The high degree of crystallinity with ordered thick lamellae crystallites obtained at a low cooling rate caused the deformation slips in the crystal blocks to be more difficult. The improvement of crystalline perfection and lamella thickening taking
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Acknowledgements This project was financially supported by the Research Grant Council of Hong Kong (CERG HKUST719/95E) and the Hong Kong University of Science & Technology (HKUST) Area of Excellence Grant (AoE97/98.EG14). Experiments were conducted with the technical supports from Advanced Engineering Materials Facility and Materials Characterisation & Preparation Facilities of HKUST. Part of the paper was presented at Third International Conference on Fracture and Strength of Solids, Hong Kong, 1997, and First Asian–Australasian Conference on Composite Materials (ACCM-1), Osaka, Japan, 1998.
References Fig. 17. Scanning electron micrographs of fracture surface and schematic fibre–matrix interaction of carbon fibre/PEEK composites cooled at: (a) 1⬚C/min; and (b) 1800⬚C/min.
place in a slow cooling process promoted the post-yield hardening behaviour as reflected by the absence of load drop after plastic yielding. The interphase bond strength decreased with increasing cooling rate through its effects on crystallinty and lamella morphology of interphase. The crystallinity in the matrix controlled the mechanical properties of bulk matrix, including the ductility, tensile strength and elastic modulus, while the lamella morphology on the fibre surface affected the interface adhesion mechanisms. The interface failure mechanisms were also a direct reflection of the interface adhesion mechanisms. The closely packed crystalline lamella interphase that was adsorbed more strongly onto the fibre surface in slow cooled specimens gave rise to a high IFSS with relatively brittle interface debonding. In contrast, the amorphous PEEK-rich interphase introduced in fast cooled specimens resulted in a low IFSS and a relatively ductile interface failure, as manifested by the thick permanently deformed matrix layer surrounding the pulled out fibre socket. The interphase adhesion represented by IFSS and ILSS, and the bulk mechanical properties of matrix, such as ductility, strength and modulus, all exhibited functionally similar variations with cooling rate. The sharp transition of these properties at about 600⬚C/min may be attributed to a transition of failure mechanisms associated with the degree of crystallinity. The sensitivity of matrix crystallinity on cooling rate was much higher than the changes in IFSS and ILSS. In summary, it is demonstrated in this study that the fibre–matrix interphase adhesion and the bulk mechanical properties of the thermoplastic composite can be optimised, if not totally maximised, by controlling the crystallinity which is affected by processing conditions, such as the cooling rate.
[1] Lustiger A, Uralil FS, Newaz GM. Processing and structural optimization of PEEK composites. Polym Compos 1990;11:65–75. [2] Saiello S, Kenny J, Nicolais L. Interface morphology of carbon fibre/ PEEK composites. J Mater Sci 1990;25:3493–6. [3] Zhang M, Xu J, Zhang Z, Zeng H, Xiong X. Effect of transcrystallinity on tensile behaviour of discontinuous carbon fibre reinforced semicrystalline thermoplastic composites. Polymer 1996;37:5151–8. [4] Kim JK, Mai YW. Interfaces in composites. In: Chou TW, editor. Materials science and technology, structure and properties in composites, 13. Weinheim, Germany: VCH, 1993. pp. 239–89. [5] Youssef Y, Denault J. Thermoformed glass fiber reinforced polypropylene: microstructure, mechanical properties and residual stresses. Polym Compos 1998;19:301–9. [6] Waterbury MC, Drzal LT. Interfacial shear strengths of carbon fibres in bisphenol-A polycarbonate. In: Ishida H, editor. Controlled interphases in composite materials, New York: Elsevier, 1990. pp. 731–9. [7] Nielsen AS, Pyrz R. Study of the influence of thermal history on the load transfer efficiency and fibre failure in carbon/polypropylene microcomposites using Raman spetroscopy. In: Proceedings of Seventh Int Conf Compos Interfaces (ICCI-VII), Shonan, Japan, 1998, paper SXI-2. [8] Schulz E, Kalinka G, Auersch W. Effect of transcrystallization in carbon fiber reinforced polyphenylene sulfide composites on the interfacial shear strength investigated with the single fiber pull-out test. J Macromol Sci: Phys B 1996;35:527–46. [9] Ye L, Scheuring T, Friedrich K. Matrix morphology and fibre pull-out strength of T700/PPS and T700/PET thermoplastic composites. J Mater Sci 1995;30:4761–9. [10] Peron B, Lowe A, Baillie C. The effect of transcrystallinity on the interfacial characteristics of polypropylene/alumina single fibre composites. Composites Part A 1996;27:839–45. [11] Marom G. The microstructure of crystalline interfaces and effects on properties of fibre-reinforced composites. In: Proc Seventh Intern Conf Compos Interfaces (ICCI-VII), Shonan, Japan, 1998, paper SVII-1. [12] Beehag A, Ye L. Fibre/matrix adhesion in thermoplastic composite: is transcrystallinity a key? In: Proc Eleventh Int Conf Compos Mater (ICCM-11), Gold Coast, Australia, 1997;4:723–730. [13] Zou YL, Netravali AN. Ethylene/ammonia plasma polymer deposition for controlled adhesion of graphite fibers to PEEK. II. Effect on fiber and fiber/matrix interface. J Adhes Sci Technol 1995;9:1505– 20. [14] Kobayashi H, Hayakawa E, Kikutani T, Takaku A. Effect of quenching and annealing on fiber pull-out from crystalline polymer matrices. Adv Compos Mater 1991;1:155–68. [15] Folkes MJ, Kalay G, Ankara A. The effect of heat treatment on the properties of PEEK and APC2. Compos Sci Technol 1993;46:77–83.
530
S.-L. Gao, J.-K. Kim / Composites: Part A 31 (2000) 517–530
[16] Gao SL, Kim JK. Interphase morphology and fibre pull-out behaviour of carbon fibre-PEEK composites. Key Engng Mater 1998;145149:811–6. [17] Gao SL, Kim JK. Effect of cooling rate on impact damage resistance of carbon fibre-PEEK laminates. In: Proc Eighth European Conf Compos Mater (ECCM-8), Naples, Italy, 1998:95–104. [18] Gao SL, Kim JK. Effect of cooling rate on interphase properties and mechanical response of carbon fibre/PEEK composites. Mater Sci Res Int 1999;5:157–62. [19] Blundell DJ, Osborn BN. The morphology of poly(aryl-ether-etherketone). Polymer 1983;24:953–8. [20] Kim JK, Zhou LM, Mai YW. Techniques for studying composite interfaces. In: Cheremisinoff NP, editor. Handbook of advanced materials testing, New York: Marcel Dekker, 1995. pp. 327–65. [21] Ohsawa T, Nakyama A, Miwa M, Hasegawa A. Temperature dependence of the critical fibre length for the glass fibre reinforced thermosetting resins. J Appl Polym Sci 1978;22:3203–17. [22] Drzal LT. Fibre–matrix interphase structure and its effect on adhesion and composite mechanical properties. In: Ishida H, editor. Controlled interphases in composite materials, New York: Elsevier, 1990. pp. 309–20. [23] Vu-Khanh T, Denault J. Effect of molding parameters on the interfacial strength in PEEK/carbon composites. J Reinf Plast Compos 1993;12:916–31. [24] Blundell DJ, Chalmers JM, Mackenzie MW, Gaskin WF. Crystalline morphology of the matrix of PEEK-carbon fibre aromatic polymer composites. I. Assessment of crystallinity. SAMPE Quart 1985;16:22–30. [25] Chen M, Chung CT. Crystallinity of isothermally and nonisothermally crystallized poly(ether ether ketone) composites. Polym Compos 1998;19:689–97. [26] Xiao XR, Hoa SV. Effect of melting history on the crystalline characteristics of polyetheretherketone aromatic polymer composite. Theor Appl Fract Mech 1990;14:49–56. [27] Davies P, Cantwell WJ, Jar PY, Richard H, Neville DJ, Kausch HH. Cooling rate effects in carbon fibre/PEEK composites. In: O’Brien TK, editor. Composite materials: fatigue and fracture, ASTM STP 1110, Philadelphia, 1991. pp. 70–88. [28] Talbott MF, Springer GS, Berglund LA. The effects of crystallinity on the mechanical properties of PEEK polymer and graphite fiber reinforced PEEK. J Compos Mater 1987;21:1056–81. [29] Velisaris CN, Seferis JC. Crystallization kinetics of polyetheretherketone (PEEK) matrices. Polym Engng Sci 1986;26:1574–81. [30] Unger WJ, Hansen JS. The effect of thermal processing on residual strain development in unidirectional graphite fibre reinforced PEEK. J Compos Mater 1993;27:59–82. [31] Yasuniwa M, Nakafuku C. Polymer 1980;12:105. [32] Partridge IK, Davies P, Parker DS, Yee AF. Yield and fracture in PES and PEEK matrix polymers and their composites. In: Proc Int Conf Polym Compos, Cranfield, UK: Plastics and Rubber Inst, 1987;5:1– 12.
[33] Uralil FS, Newaz GM, Lustiger A. Processing effects and damage tolerance in poly(etheretherketone) composites. Polym Compos 1992;13:7–14. [34] Lee Y, Porter RS. Crystallization of polyetheretherketone (PEEK) in carbon fibre composites. Polym Engng Sci 1986;26:633–9. [35] Waddon AJ, Hill MJ, Keller A, Blundell DJ. On the crystal texture of linear polyaryls (PEEK, PEK and PPS). J Mater Sci 1987;22:1773– 84. [36] Chen EJH, Hsiao BS. The effects of transcrystalline interphase in advanced polymer composites. Polym Engng Sci 1992;32:280–6. [37] Deporter J, Baird DG. The effects of thermal history on the structure/ property relationship in polyphenesulfide/carbon fibre composites. Polym Compos 1993;14:201–13. [38] Tregub A, Harel H, Marom G, Migliaresi C. The influence of thermal history on the mechanical properties of poly(ether ether ketone) matrix composite materials. Compos Sci Technol 1993;48:185–90. [39] Jar PY, Kausch HH. Annealing effect on mechanical behaviour of PEEK. J Polym Sci Polym Phys 1992;30:775–8. [40] Weibull W. A statistical distribution function of wide applicability. J Appl Mech 1951;18:293–6. [41] Galiotis C, Young RJ, Yeung PHJ, Batchelder DN. The study of model polydiacetylene/epoxy composites. I. The axial strain in the fibre. J Mater Sci 1984;19:3640–8. [42] Asloun EM, Nardin M, Schultz J. Stress transfer in single-fibre composites: effect of adhesion, elastic modulus of fibre and matrix, and polymer chain mobility. J Mater Sci 1989;24:1835–44. [43] Ogata N, Yasumoto H, Yamasaki K, Hu Y, Ogihara T, Yanagawa T, Yoshida K, Yamada Y. Evaluation of interfacial properties between carbon fibres and semicrystalline thermoplastic matrices in singlefibre composites. J Mater Sci 1992;27:5108–12. [44] Cox HL. The elasticity and strength of paper and other fibreous materials. Brit J Appl Phys 1952;3:72–79. [45] Drzal LT, Rich MJ, Koenig MF, Lloyd PF. Adhesion of graphite fibres to epoxy matrices: II. Effect of fibre finish. J Adhes 1983;16:133–52. [46] Pisanova EV, Zhandarov SF. On the mechanism of failure in microcomposites consisting of single glass fibres in a thermoplastic matrix. Compos Sci Technol 1997;57:937–43. [47] Kim JK, Mai YW. Engineered interfaces in fibre reinforced composites. Oxford, UK: Elsevier, 1998 p. 43–92. [48] Brady RL, Porter RS. Interfacial adsorption and crystallization of polycarbonate in carbon fibre composites. J Appl Polym Sci 1990;39:1873–85. [49] Felix JM, Gatenholm P. Effect of transcrystalline morphology on interfacial adhesion in cellulose/polypropylene composites. J Mater Sci 1994;29:3043–9. [50] Clark RLJ, Kander RG, Sauer BB. Nylon 66/poly(vinyl pyrrolidone) reinforced composites. I. Interphase microstructure and evaluation of fiber–matrix adhesion. Composites Part A 1999;30:27–36.