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Damage development during flexural fatigue of carbon fibre-reinforced PEEK
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Composites 26 (1995) 355-370
U T T E R W O R T H E I N E M A N N
Elsevier Science Limited Printed in Great Britain. 0010-4361/95/$10.00
Damage development during flexural fatigue of carbon fibre-reinforced PEEK G. Dillon* Laboratory for Manufacturing and Productivity, Massachusetts Institute of Technology, Cambridge, MA 02139, USA and M. Buggy Department of Materials Engineering and Industrial Chemistry, University of Limerick, Castletroy, Limerick, Ireland Fatigue damage growth mechanisms in [0116and [0, 9014slaminates of APC-2 subjected to three-point flexure were studied using scanning electron microscopy. Building on an understanding gained from the study of statically tested samples, specificfatigue mechanisms were identified. For each laminate geometry a coherent progression of damage build-up is proposed. (Keywords: flexural fatigue; damage mechanisms; APC-2 laminates; scanning electron microscopy)
INTRODUCTION It has long been recognised that there is a need for a fundamental technique for analysing and understanding the large volume of fatigue data available for composite materials 1. Previous studies of the fatigue behaviour of poly(ether ether ketone) (PEEK) matrix composites have concentrated on establishing the fatigue performance of this material 2'3. This paper is concerned with the use of scanning electron microscopy in the determination of fatigue damage growth mechanisms in AS-4 carbon fibre-reinforced PEEK (APC-2) laminates. The object of this work is to add to the general understanding of the fatigue behaviour of APC-2 by coherently describing the sequence of events leading to the failure of laminates of various geometries subjected to dynamic flexural stressing. Samples of [0]16 and [0, 9014s APC-2 laminates were prepared by compression moulding following recommended procedures. Their static mechanical properties and flexural fatigue behaviour were measured in threepoint bending using a jig conforming to BS 2782 (3) Method 304E. Details of sample preparation, numerical results and fatigue damage modelling are reported elsewhere4. In all cases the plane of the failure surface was essentially an extension of the central plane of the loading anvil. The important features of composites fractography have been covered in detail elsewhere5-9. In this study the technique employed for the identification of fatigue damage development mechanisms was scanning electron microscopy. Fracture surfaces of both statically and dynamically tested laminates of APC-2 were examined * To w h o m correspondence should be addressed
in an effort to establish the sequence of events leading to final failure. The discussion will focus on unidirec= tional and cross-plied laminates as the attempts on fractographic analysis of angle-plied laminates proved difficult, due to the complexity of the fracture surfaces. The scanning electron microscopy was carried out on a Jeol JSM80 instrument. Samples were gold coated to a depth of 45 nm and examined using an accelerating voltage of 20 kV. DAMAGE DEVELOPMENT IN STATIC FAILURES To establish the basic fractographic features and to verify that they were similar to those previously reported 7-9, fractographs of samples subjected to static flexural testing were examined. Unidirectional laminates A low magnification view of a fracture surface generated in a static flexural test of a unidirectional laminate is shown in Figure 1. The fracture surface can be divided into tensile-dominated failure (at the top of the figure) and compressive failure. The changeover is clearly visible as an almost perfectly straight line at the neutral axis. This is typical of flexural failures of carbon fibre-reinforced plastic (CFRP) laminates as reported by other investigators 9-11. It has been established that the relative areas of tensile and compressive failure on flexural fracture surfaces can give an indication of the initiation mechanism1~ The changeover from tensile to compressive failure (hereafter referred to as the changeover line or neutral axis) shown in Figure I indicates that crack initiation occurred almost simultaneously at both faces.
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Figure 1 Static flexural failure of [0116 specimen (X47)
Figure 2 Static failure of [0116specimen, tensile side (x65)
The smoothness of the compressive failure surface in Figure 1 suggests that crack propagation was very rapid, as no evidence of crack slowing was found. Also, delamination was not noted on either the tensile or compressive fracture surface. An angled view of the tensile fracture surface is shown in Figure 2. This shows the variation in the degree of fibre pull-out from the tensile face to the changeover line. Directly attributable fibre failure 5 is rare, as fibre clusters tend to be small and very dispersed. This suggests the random initiation of fibre failures on the tensile face. As the failure progresses into the sample, fibre stresses tend towards pure flexure due to the increasing radius of curvature. Figures 3 and 4 illustrate features of brittle and ductile matrix failures respectively which have occurred as a result of different crack propagation rates. Figure 3 shows fracture occurring close to the tensile
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face. The brittle appearance of the matrix adhering to the fibres indicates a rapid crack growth rate. Figure 4 shows the failure occurring very close to the changeover line. Ductile drawing has occurred to such an extent that fibres have been completely stripped of matrix material. In view of the above, the following fracture sequence for unidirectional APC-2 laminates subjected to static three-point flexure is proposed. Crack initiation occurs simultaneously at both sample faces. On the tensile side, random fibre breakage occurs with initially rapid crack propagation rates. Crack propagation slows down as tensile failure approaches the neutral axis, leading to varying levels of fibre pull-out. On the compressive side the crack progresses at a uniform rapid rate. Final failure occurs when tensile and compressive cracks meet at the neutral axis.
Damage development during flexural fatigue: G. Dillon and M. Buggy
Figure 3 Static failure of [0116 specimen, showing brittle failure near the tensile surface (x 1300)
Figure 4 Static failure of [0116 specimen, showing ductile failure near the neutral axis (x 1300)
Cross-plied laminates A low magnification view of a statically fractured cross-plied laminate is shown in Figure 5. If plies are numbered 1 to 16 starting at the compressive face (as illustrated in Figure 8), then it is clear that the changeover line occurs close to the boundary between plies 3 and 4. The remainder of the fracture surface exhibits tensile characteristics. Figure 6 shows an angled view of the tensile failure surface where a complex crack pattern exists, and extensive delamination has occurred. Large sections of transverse plies are pulled off the fracture surface, and, as shown in Figure 7, transverse ply cracking occurs on a number of different planes within a 90 ~ ply. It should be borne in mind, however, that a lot of this transverse ply cracking will have occurred after the fracture has reached an unstable phase. Nonetheless, many transverse ply cracks not associated with postfailure events were found, suggesting the development of
complex transverse ply crack patterns before failure. Again, fibre failure is seen to occur on a number of different planes, suggesting that fibre failure was spatially random in the 0~ plies. However, fibre ends in areas close to the ply interfaces showed radial patterns, suggesting crack propagation away from a cracked 90 ~ ply. Such areas were generally narrow and a good example is shown close to the centre line of Figure 6 at A-A. This suggests that stress intensification due to transverse ply cracking led to fibre breakage at the interface. Figure 8 shows crack propagation between plies 14 and 15. The ductile drawing indicates slow crack propagation and the figure also shows where fibres lay before crack propagation occurred. An important feature is the tensile nature of matrix failure between these plies. Shear abrasion of mating surfaces was not found on statically tested samples, indicating that delamination must generally have occurred towards the end of a load sequence.
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Figure 5 Static failure of [0, 9014sspecimen (x43)
Figure 6 Static failure of [0, 9014sspecimen, showing complex crack pattern (x75)
Figure 7 Static failure of [0,9014s specimen, showing multiple transverse cracks (x300)
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Figure 8 Static failure of [0, 9014s specimen, showing ductile strain magnification between plies (x2000)
Figure 9 Static failure of [0, 9014~specimen, showing brittle spherulitic pattern on a transverse ply (• 3300)
Figures 9 and 10 show several identifiable features of transverse ply failure. The clear spherulitic pattern shown in Figure 9 is indicative of a fast brittle fracture. The star-like patterns are spherulites within the matrix material and the character of such features can be used to judge crack propagation rates. When crack propaga-
tion rates are rapid the spherulitic patterns tend to be readily visible, as in Figure 9. River markings as described by Purslow 5 are clearly visible in Figure 9. It was noted that there was a definite tendency for these markings to occur at ply boundaries, as shown at A in Figure 9. Their general orientation suggests crack prop-
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Figure 10 Staticfailure of [0, 9014sspecimen, showingpeel failure within ply 15 (X 1500) agation away from interfaces. As the maximum tensile strain generated within a 90 ~ ply is likely to occur at the boundary closest to the tensile face of the sample, this is not surprising. In some instances it is possible to observe minor cracks opposed to the primary crack propagation direction. If the following crack sequence proposed for plies 13, 14 and 15 is considered, this feature can be explained. A transverse ply crack initiates in ply 15, close to the ply interface nearest the tensile face of the sample. This crack grows rapidly into the ply, leading to stress intensification in the neighbouring plies. It is proposed that at this stage extensive fibre breakage has occurred in both plies 14 and 16, but both plies are capable of bearing high loads. The stress intensification leads to matrix failure in ply 13; again at the interface closest to the tensile face. This in turn causes further intensification of stress in ply 14 and eventual failure of the matrix close to the interface between plies 14 and 15. In this way, ply 15 is subjected to tensile cracking from both ply interfaces. This proposal is supported by the observation that river markings, close to the ply interfaces on either side of the same ply, were frequently oriented in different directions, confirming that transverse ply cracking initiated from a number of different sources. Figure 10 shows cusp formation and this indicated that a significant shear component of stress existed in the plane of the transverse crack, at the crack tip. The orientation of these peel cusps indicates that crack propagation was from left to right. The coarseness of their formation suggests that ply separation was significant, before they were generated. The appearance of peel cusps gives an important clue as to lateral crack growth
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mechanisms within a ply. Frequently a change in cusp orientation is noted. In some instances the point about which this change in orientation takes place can be recognised as the crack origin. Figure 11 shows a distinct change in crack propagation rate occurring in ply 8. In the region marked A, ductile drawing is significant and stripping of matrix material from fibres is clearly evident. In area B the welldefined spherulitic pattern suggests a more rapid crack propagation rate. Figure 12 shows a higher magnification of the ductile failure surface and, significantly, spherulitic features can not be noted even at this very high magnification. The above indicates a sudden slowing of crack propagation rates within the ply. A possible explanation is that transverse cracks are initiated at different positions and on different planes within this 90 ~ ply. The crack propagation directions were laterally opposed. As the crack fronts propagated past one another, a thin strip of the ply was separated from the main fracture surface. The energy of propagation was then transferred to bending of the pulled-out strip and consequently crack propagation changed from brittle matrix cracking to ductile peeling. Extremely rapid crack propagation within plies 8 and 9 is shown in Figure 13. Here, the spherulitic pattern is not visible at x 1000 magnification. However, when this surface was viewed at extremely high magnifications, evidence of the spherulitic formation was again visible. In earlier work 7 it was suggested that crack propagation rates can be categorized as slow ductile, intermediate and brittle, on the basis of spherulitic definition. It is proposed here that, in the case of brittle crack propaga-
Damage development during flexural fatigue: G. Dillon and M. Buggy
Figure 11 Static failure of [0, 9014s specimen, showing change from brittle to ductile failure in ply 8 (x 1500)
Figure 12 Static failure of [0, 9014sspecimen, showing ductile drawing
(x5ooo)
tion, an indication of crack growth rate can be estimated from the magnification necessary to resolve spherulitic features. The spherulitic features are visible because of the differential responses of the constituent crystalline and amorphous regions. Effectively, higher crack propagation rates lead to smoother surfaces requiring higher magnifications to observe this effect. Figure 13 also shows the atypical feature of interlaminar cracking between plies 8 and 9 (x-x). In general, this interface
between two 90 ~ plies remained perfectly intact and plies tended to act as one. The above discussion leads to a fairly clear picture of damage development in cross-plied laminates of statically tested APC-2 laminates. A network of transverse ply cracks builds up in the plies close to the tensile surface, initiation occurring mainly in the regions close to interply layers. Fibre breaks occur on the tensile surface and at ply interfaces, leading to intensification of stresses acting on transverse plies. Progressive fibre failures lead to increased stresses on plies closer to the neutral axis and transverse cracking penetrates deeper into the sample. Compressive cracking occurs in the first few plies and progression into the sample occurs as a straight front. The neutral axis is closer to the compressive surface than in the unidirectional case. This is due to the fact that transverse plies contribute greater support in compression than in tension. Final failure occurs when transverse cracking and fibre failure lead to stress intensity of a sufficient magnitude to fracture the remaining longitudinal fibres. Post-failure delamination and transverse ply pull-out then occur. Finally, it is worth pointing out that crack propagation rates can be accurately traced through each transverse ply. In general the fastest crack propagation rates were noted in the central plies. This was so because they act as a single, much thicker transverse ply, thereby being subject to a lesser degree of crack constraint from 0~ plies. Also, a large degree of damage had already occurred in neighbouring plies before central ply failure. Thus, failure of the central plies is associated with the final rapid failure of the test specimen.
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Figure 13 Static failure of [0, 9014sspecimen, showing interlaminar crack between plies 8 and 9 (x 1000)
Figure 14 Dynamic failure of [0116 specimen tested at 90% of static strength (•
DAMAGE DEVELOPMENT IN FATIGUE FAILURES Fracture surfaces generated by flexural fatigue of unidirectional and cross-plied laminates showed many of the identifiable features found in static failures. However, a number of features particularly associated with fatigue were identified.
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Fatigue failures of unidirectional laminates Figure 14 shows a general view of a sample subjected to a maximum fatigue stress of 90% of its ultimate flexural strength. The sample shown survived 97 load cycles to this stress level. As can be seen by comparison with Figure 1, the fracture surface is similar to that generated in a static test. However, there are two major differences.
Damage development during flexural fatigue: G. Dillon and M. Buggy
Figure 15 Dynamic failure of [0]~6specimen tested at 70% of static strength (x90)
First, the compression failure surface is more complex, particularly near the compression face of the sample. Second, the compression surface is more noticeably curved than in the static case. The complex crack pattern near the surface suggests a gradual build-up of damage by fatigue loading. The horizontal lines x-x, y-y, and z-z are associated with crack arrest points during the last few cycles of loading. The tendency towards compression surface curvature in fatigued laminates is not easily explained but was observed as a general phenomenon. In the static situation the planar character of the compression surface can be explained by the fact that crack propagation occurs in a single stage, and is very rapid. The stored elastic strain energy is released suddenly as unstable cracks initiate from both sample faces. The radius of curvature induced by bending causes cracks on the compression side to propagate at a slight angle to the plane normal to the fibre directions. Once initiated, crack propagation is so rapid that the crack direction does not change until the crack front is very close to the neutral axis. In the dynamic situation crack propagation is much more gradual and directional changes can occur. This may be due to an increasing offaxis stress component, for example shear, as the crack progresses through the compressive region. Supporting this contention is the observation that this effect was more marked in samples subjected to lower fatigue loads. The tensile surface is very similar to that of the statically tested sample. In fact, no significant difference between these two fracture surfaces was noted. However, there was a tendency towards greater fibre pull-out in the fatigued samples. This suggests a more random fibre breakage pattern in the fatigue situation. An angled view of the tensile surface of a sample subjected to a maximum fatigue stress of 70% of the ultimate flexural strength is shown in Figure 15. Again the tensile surface is similar to that of the static case. Single fibres in isolation indicate fibre failure occurring before overall final fracture and these were found to be more in evidence in the fatigued samples. Failure seems to have occurred in intraply bands, suggesting that in this case
crack growth occurred as a number of discrete steps on the tensile side. A number of fibre bundles were found near the tensile face with adhering matrix material, showing ductile shearing. This may indicate either a slowing of crack propagation rates in these regions or, more likely, shear abrasion of matrix material after individual fibre failures had occurred. Clear evidence of abrasion following fibre failure can be seen in Figure 16. Bare fibres were found in a region where crack propagation was rapid. However, as pointed out previously, bare fibres are normally associated with ductile failure. It is considered that a longitudinal crack propagated in a brittle fashion through the area illustrated. Fatigue abrasion of the mating surfaces led to removal of the matrix material surrounding the fibre. As this feature was not noted in statically tested samples, it is suggested that an indication of the amount of longitudinal matrix cracking that has occurred can be estimated from the level of fatigue abrasion observed. As with statically tested samples, the ductility of matrix failure increased towards the neutral axis. As illustrated clearly in Figure 17, fibre damage is much more localized, which may be due to the decreasing radius of curvature. It is also possible that high shear stresses are generated here, as fibre fracture seems to have occurred at 45 ~ to the normal plane in many cases. It is most likely, however, that the fibre stresses tended towards pure flexure as the tensile failures approached the neutral axis. Figure 18 shows a general view of the compression surface of the above sample. This surface can be divided into three distinct regions, in terms of smoothness. Higher magnification views show that the variation in surface smoothness is due to different levels of fatigue abrasion. Figure 19 shows what can be termed low abrasion. Some evidence of the fibrous nature of the material is still visible in this area. In the top left at A, for instance, the general outline of a fibre end can be seen. Figure 20 shows medium level abrasion. Here, the underlying fibres and matrix material have been abraded to such an extent that the individual phases can no longer be distinguished. The entire surface in this region is
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Figure 16 Dynamic failure of [0116specimen tested at 70% of static strength, showing bare fibres indicating fatigue abrasion (x 6000)
Figure 17 Dynamic failure of [0116 specimen tested at 70% of flexural strength, showing fibre breakage near neutral axis (x 1800)
Figure 18 Dynamic failure of [0116 specimen, showing different levels of compressive abrasion (x330)
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Figure 19 Dynamic failure of [0]16 specimen tested at 70% of flexural strength, showing low level abrasion (x 3000)
Figure 20 Dynamic failure of [0116specimen tested at 70% of flexural strength, showing medium level abrasion at interface between plies 2 and 3 (x3300)
covered with damage debris. Figure 21 shows high level abrasion where damage debris has been smoothed off the fracture surface. A series of horizontal lines can also be seen in this fractograph, indicating gradual growth of the compression crack through this region. The dispersion of low, medium and high level abrasion regions throughout the fracture surface was quite complex. In general, a region of medium abrasion is seen close to the compressive surface. This is associated with the relatively rapid initial propagation of a compression crack. The growth rate then slows into the high abrasion region. The smoothness of this region indicates a slow crack propagation rate, suggesting that the greater part of fatigue life is spent in this slow propagation phase. Rapid crack propagation towards the end of fatigue life gives rise to the low level abrasion regions found close to the neutral axis. Finally, a region of pure flexural failure (i.e. no abrasion) occurs at or near the neutral axis. This is associated with the final cycle of testing. It should be noted that regions of high abrasion do not extend across the entire specimen width. Typically, three such regions occur in any one sample, generally
bounded by regions of medium and low abrasion. The obvious implication is that a compression crack is initiated across the sample width and splits into three or four growing cracks, which gradually progress into the sample. Initial cracks grow from the induced fibre damage created by the central loading anvil. In some instances an increase in crack propagation rate occurs in a region of otherwise slow crack growth rate. This is generally accompanied by delamination, as shown in Figure 18 at x-x. Figure 22 shows the shear character of a crack that propagated longitudinally within a ply near the surface. It is thought that growing compression cracks can lead to an increase in matrix shear stresses at the crack tip. The crack grows from the compression surface until encountering an interface. Crack tip shear stresses can then cause the crack to propagate interply. Effectively, an edge notch is introduced to each half of the specimen at the centre and the sample acts similarly to two edge-notch flexure samples. This causes a temporary increase in the rate of crack growth. Previous studies of different crack growth rates in fatigue of composites were not found in the literature. Compressive fatigue fracture has rarely been the subject COMPOSITES Volume 26 Number 5 1995
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Figure 21 Dynamic failure of [0116specimen tested at 70% of flexural strength, showing high level abrasion (x3300)
Figure 22 Dynamic failure of [0116 specimen tested at 70% of flexural strength, showing intraply shear failure near tensile surface (•
of fractographic studies presumably because it has generally been assumed that post-failure abrasion renders fracture surfaces uninterpretable. On the contrary, it is believed that this abrasion provides a useful clue to propagation mechanisms and rates. Interestingly, different abrasion levels were not noted in comparative carbon fibre/epoxy laminates (T300/Cycom 985GT 6095 III CFRP of geometries [0]20, [9015s and [+4515s) studied in the present work. This suggests that progressive damage mechanisms in carbon fibre/epoxy systems are associated with gradual build-up of sub-surface debonds and matrix cracks, whereas in the carbon fibre/thermoplastic system gradual growth of compression cracks represents the main recognisable progressive damage mechanism. Fatigue damage development in unidirectional laminates subjected to low fatigue stresses can be summarized as follows. A compressive crack is induced by load
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member stress intensification. Random fibre breaks occur on the tensile surface. Bands of tensile failure progress into the sample, and a high degree of shear abrasion of failed regions occurs. A number of slowly propagating compression cracks propagate from the original crack. These cracks grow gradually into the sample, giving rise to smooth abraded surfaces. Occasional delaminations occur where high shear stresses are generated at compression crack tips. Final failure occurs when the depth of compression cracking causes fibre stress intensification of a sufficient magnitude to initiate unstable fracture. The degree and type of abrasion can be used to estimate the approximate number of cycles, and consequently the applied stress. Medium to high abrasion is associated with fatigue stresses in the region of 70 to 75% of ultimate strength, whereas low abrasion levels are associated with higher stresses.
Damage development during flexural fatigue: G. Dillon and M. Buggy
Figure 23 Dynamic failure of [0, 9014sspecimen tested at 90% of flexural strength (x43)
Figure 24 Dynamic failure of [0, 9014sspecimen tested at 80% of flexural strength, showing lateral crack growth (x 1000)
Finally, it should be noted that migration of the neutral axis was not observed in unidirectional samples in this study.
Fatigue failures of cross-plied laminates Figure 23 shows a general view of a cross-plied sample subjected to fatigue at 90% of ultimate flexural strength. Comparison with Figure 15 shows that, on a low magnification level, tensile fracture surfaces are similar to those generated in the static situation. In fact, very few differ-
ences between high stress fatigue damaged laminates and statically tested laminates were noted. The degree of delamination in both cases was similar and the same tendency towards crack propagation away from ply boundaries was noted. Figure 24 shows a feature that is believed to be particular to fatigued samples. The feature is a rib, identified 7 as an indication of a change in crack propagation rate. However, though such changes were noted in statically tested samples, ribs of the kind illustrated were not. Closer examination of matrix material
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Figure 25 Dynamic failure of [0, 9014sspecimen tested at 90% of flexural strength, showing striations mirroring underlying fibre surface profile (x20 000)
on either side of the rib shows that crack propagation was in both cases rapid. It is proposed therefore that the rib represents a transverse crack arrest line, between load cycles. The general curvature of this line again suggests that initiation occurred at a ply interface. Features in Figure 25 were initially believed to be fatigue striations. However, further examination showed that these striations simply matched those already present on fibre surfaces caused by the fibre production process. Figure 26 shows a general view of a cross-plied specimen subjected to fatigue at 75% of ultimate flexural strength. As can be seen, the changeover line occurs in the middle of ply 2, suggesting that eventual failure was tensile-dominated. The degree of delamination and fibre pull-out is much greater than in statically tested and high stress fatigue samples. This indicates that gradual buildup of fibre breaks and matrix cracks characterized the
earlier part of the damage process. Towards the end of cycling delamination occurs, leading to stress intensification in surviving fibres. Different crack propagation rates were noted within transverse plies. Figure 27 shows brittle failure 7 in a ply in which extremely rapid crack propagation rates were also noted. Compression cracks generated at the compression sample faces showed some evidence of different levels of abrasion. However, it is believed that surface compression cracks immediately change direction on meeting a ply boundary. A significant amount of delamination in surface plies was noted on the compression side. Also, there was clear evidence of intraply splitting in longitudinal plies. This suggests gradual yielding of matrix material, allowing development of shear cracks within 0 ~ plies. On the tensile side a number of longitudinal matrix cracks were seen to run parallel to the fibres in the 0 ~ plies. Constant-moment flexure tends to induce anticlastic curvature. This curvature is prevented by the presence of the load member in three-point bend flexure, and consequently transverse stresses are induced in the surface plies. These stresses lead to a gradual build-up of tensile surface cracks during fatigue. Evidence was found that suggested that this type of cracking also occurs in longitudinal plies beneath the specimen surface. The appearance of spherulitic features, shown in Figure 28, was particular to fatigued samples. This is very similar to the pattern noted in the interply regions. Both surfaces showed a definite ductile tensile character. The origin of this pattern is not clearly understood. It may be due to gradual progression of a fatigue crack through a spherulite, where fibrils are drawn during successive load cycles. It is also possible that this pattern is associated with a crack propagation rate not previously noted. The fatigue damage development process in crossplied laminates can be summarized as follows. Compression cracking occurs on the sample face in the first few cycles. Sub-surface transverse ply cracking initiates, primarily, near ply interfaces. Fibre breakage occurs in the outermost tensile plies. Intraply longitudinal
Figure 26 Dynamic failure of [0, 9014sspecimen tested at 75% of flexural strength. Angled view at 30~ to show crack route near compressive surface (x47)
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Damage development during flexural fatigue: G. Dillon and M. Buggy CONCLUDING REMARKS
Figure 27 Dynamic failure of [0, 9014s specimen tested at 75% of flexural strength, showing intermediate brittle failure on pulled-out 90~ ply
(x4800)
matrix cracks are initiated by high interfacial shear stresses parallel to the ply interfaces. Longitudinal matrix cracks are also initiated by high transverse stresses generated in the 0 ~ plies. Growing compression cracks and transverse ply cracks penetrate interply regions giving rise to delaminations. All the above mechanisms continue to operate until unstable cracking occurs in the final cycle.
The present study illustrates the usefulness of scanning electron microscopy in identifying failure mechanisms in composite laminates. Previous investigators have shown that a good deal of useful information on static failure mechanisms can be attained using this technique. This work shows that many of the features identified in the static situation can also be seen in samples subjected to flexural fatigue loading and that additional features particular to the dynamic loading situation can be recognised. Fatigue striations were not seen in the form normally associated with polymeric materials, except on a coarse scale at very high stress levels, and then only in unidirectional samples. More useful information on the fatigue history of a laminate can be gained by studying the level of fatigue abrasion in compression cracks generated directly beneath the load member. Higher levels of abrasion indicate longer lives and therefore lower stress levels, in the region of 70% of the ultimate strength of the laminate. An indication of the level of longitudinal ply matrix cracking occurring can be estimated from the degree of matrix fatigue abrasion noted. Cross-plied laminates showed extensive transverse and longitudinal ply cracking and interlaminar shearing. The usefulness of the technique is somewhat limited by laminate geometry, samples containing angled plies showing extremely complex shear failures, although shearinitiated progressive compression failure was identified. Nonetheless, it is felt that more attention should be focused on scanning electron microscopy as a means of identifying failure mechanisms in composite materials.
Figure 28 Dynamic failure of [0, 9014s specimen tested at 75% of flexural strength, showing coarse spherulitic fracture on 90~ ply (x4800)
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Damage development during flexural fatigue: G. Dillon and M. Buggy REFERENCES 1 Talreja, R. 'Fatigue of Composite Materials', Technomic Publishing Co. Inc., Westport, CT, 1987 2 Hartness, J.T. in 'Proc. 28th Nat. SAMPE Syrup.' 12-14 April 1983, p. 535 3 Kim, R.Y. and Hartness, J.T. in 'Proc. 29th Nat, SAMPE Syrup.' 3-5 April 1984, p. 765
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Buggy, M. and Dillon, G. Composites 1991, 22 (3), 191 Purslow, D. Composites 1981, 12 (4), 241 Purslow, D. Composites 1986, 17 (4), 289 Purslow, D. Composites 1987, 18 (6), 365 Purslow, D. Composites 1988, 19 (2), 115 Purlsow, D. Composites 1988, 19 (5), 358 Croman, R.M. in 'Proc. ICCM and ECCM', London, UK, July 1987, p. 4.76 11 Bader, M.G. and Johnson, M. Composites 1974, 5, 58 4 5 6 7 8 9 10
Composites 26 (1995) 355-370
U T T E R W O R T H E I N E M A N N
Elsevier Science Limited Printed in Great Britain. 0010-4361/95/$10.00
Damage development during flexural fatigue of carbon fibre-reinforced PEEK G. Dillon* Laboratory for Manufacturing and Productivity, Massachusetts Institute of Technology, Cambridge, MA 02139, USA and M. Buggy Department of Materials Engineering and Industrial Chemistry, University of Limerick, Castletroy, Limerick, Ireland Fatigue damage growth mechanisms in [0116and [0, 9014slaminates of APC-2 subjected to three-point flexure were studied using scanning electron microscopy. Building on an understanding gained from the study of statically tested samples, specificfatigue mechanisms were identified. For each laminate geometry a coherent progression of damage build-up is proposed. (Keywords: flexural fatigue; damage mechanisms; APC-2 laminates; scanning electron microscopy)
INTRODUCTION It has long been recognised that there is a need for a fundamental technique for analysing and understanding the large volume of fatigue data available for composite materials 1. Previous studies of the fatigue behaviour of poly(ether ether ketone) (PEEK) matrix composites have concentrated on establishing the fatigue performance of this material 2'3. This paper is concerned with the use of scanning electron microscopy in the determination of fatigue damage growth mechanisms in AS-4 carbon fibre-reinforced PEEK (APC-2) laminates. The object of this work is to add to the general understanding of the fatigue behaviour of APC-2 by coherently describing the sequence of events leading to the failure of laminates of various geometries subjected to dynamic flexural stressing. Samples of [0]16 and [0, 9014s APC-2 laminates were prepared by compression moulding following recommended procedures. Their static mechanical properties and flexural fatigue behaviour were measured in threepoint bending using a jig conforming to BS 2782 (3) Method 304E. Details of sample preparation, numerical results and fatigue damage modelling are reported elsewhere4. In all cases the plane of the failure surface was essentially an extension of the central plane of the loading anvil. The important features of composites fractography have been covered in detail elsewhere5-9. In this study the technique employed for the identification of fatigue damage development mechanisms was scanning electron microscopy. Fracture surfaces of both statically and dynamically tested laminates of APC-2 were examined * To w h o m correspondence should be addressed
in an effort to establish the sequence of events leading to final failure. The discussion will focus on unidirec= tional and cross-plied laminates as the attempts on fractographic analysis of angle-plied laminates proved difficult, due to the complexity of the fracture surfaces. The scanning electron microscopy was carried out on a Jeol JSM80 instrument. Samples were gold coated to a depth of 45 nm and examined using an accelerating voltage of 20 kV. DAMAGE DEVELOPMENT IN STATIC FAILURES To establish the basic fractographic features and to verify that they were similar to those previously reported 7-9, fractographs of samples subjected to static flexural testing were examined. Unidirectional laminates A low magnification view of a fracture surface generated in a static flexural test of a unidirectional laminate is shown in Figure 1. The fracture surface can be divided into tensile-dominated failure (at the top of the figure) and compressive failure. The changeover is clearly visible as an almost perfectly straight line at the neutral axis. This is typical of flexural failures of carbon fibre-reinforced plastic (CFRP) laminates as reported by other investigators 9-11. It has been established that the relative areas of tensile and compressive failure on flexural fracture surfaces can give an indication of the initiation mechanism1~ The changeover from tensile to compressive failure (hereafter referred to as the changeover line or neutral axis) shown in Figure I indicates that crack initiation occurred almost simultaneously at both faces.
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Figure 1 Static flexural failure of [0116 specimen (X47)
Figure 2 Static failure of [0116specimen, tensile side (x65)
The smoothness of the compressive failure surface in Figure 1 suggests that crack propagation was very rapid, as no evidence of crack slowing was found. Also, delamination was not noted on either the tensile or compressive fracture surface. An angled view of the tensile fracture surface is shown in Figure 2. This shows the variation in the degree of fibre pull-out from the tensile face to the changeover line. Directly attributable fibre failure 5 is rare, as fibre clusters tend to be small and very dispersed. This suggests the random initiation of fibre failures on the tensile face. As the failure progresses into the sample, fibre stresses tend towards pure flexure due to the increasing radius of curvature. Figures 3 and 4 illustrate features of brittle and ductile matrix failures respectively which have occurred as a result of different crack propagation rates. Figure 3 shows fracture occurring close to the tensile
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face. The brittle appearance of the matrix adhering to the fibres indicates a rapid crack growth rate. Figure 4 shows the failure occurring very close to the changeover line. Ductile drawing has occurred to such an extent that fibres have been completely stripped of matrix material. In view of the above, the following fracture sequence for unidirectional APC-2 laminates subjected to static three-point flexure is proposed. Crack initiation occurs simultaneously at both sample faces. On the tensile side, random fibre breakage occurs with initially rapid crack propagation rates. Crack propagation slows down as tensile failure approaches the neutral axis, leading to varying levels of fibre pull-out. On the compressive side the crack progresses at a uniform rapid rate. Final failure occurs when tensile and compressive cracks meet at the neutral axis.
Damage development during flexural fatigue: G. Dillon and M. Buggy
Figure 3 Static failure of [0116 specimen, showing brittle failure near the tensile surface (x 1300)
Figure 4 Static failure of [0116 specimen, showing ductile failure near the neutral axis (x 1300)
Cross-plied laminates A low magnification view of a statically fractured cross-plied laminate is shown in Figure 5. If plies are numbered 1 to 16 starting at the compressive face (as illustrated in Figure 8), then it is clear that the changeover line occurs close to the boundary between plies 3 and 4. The remainder of the fracture surface exhibits tensile characteristics. Figure 6 shows an angled view of the tensile failure surface where a complex crack pattern exists, and extensive delamination has occurred. Large sections of transverse plies are pulled off the fracture surface, and, as shown in Figure 7, transverse ply cracking occurs on a number of different planes within a 90 ~ ply. It should be borne in mind, however, that a lot of this transverse ply cracking will have occurred after the fracture has reached an unstable phase. Nonetheless, many transverse ply cracks not associated with postfailure events were found, suggesting the development of
complex transverse ply crack patterns before failure. Again, fibre failure is seen to occur on a number of different planes, suggesting that fibre failure was spatially random in the 0~ plies. However, fibre ends in areas close to the ply interfaces showed radial patterns, suggesting crack propagation away from a cracked 90 ~ ply. Such areas were generally narrow and a good example is shown close to the centre line of Figure 6 at A-A. This suggests that stress intensification due to transverse ply cracking led to fibre breakage at the interface. Figure 8 shows crack propagation between plies 14 and 15. The ductile drawing indicates slow crack propagation and the figure also shows where fibres lay before crack propagation occurred. An important feature is the tensile nature of matrix failure between these plies. Shear abrasion of mating surfaces was not found on statically tested samples, indicating that delamination must generally have occurred towards the end of a load sequence.
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Figure 5 Static failure of [0, 9014sspecimen (x43)
Figure 6 Static failure of [0, 9014sspecimen, showing complex crack pattern (x75)
Figure 7 Static failure of [0,9014s specimen, showing multiple transverse cracks (x300)
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Figure 8 Static failure of [0, 9014s specimen, showing ductile strain magnification between plies (x2000)
Figure 9 Static failure of [0, 9014~specimen, showing brittle spherulitic pattern on a transverse ply (• 3300)
Figures 9 and 10 show several identifiable features of transverse ply failure. The clear spherulitic pattern shown in Figure 9 is indicative of a fast brittle fracture. The star-like patterns are spherulites within the matrix material and the character of such features can be used to judge crack propagation rates. When crack propaga-
tion rates are rapid the spherulitic patterns tend to be readily visible, as in Figure 9. River markings as described by Purslow 5 are clearly visible in Figure 9. It was noted that there was a definite tendency for these markings to occur at ply boundaries, as shown at A in Figure 9. Their general orientation suggests crack prop-
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Figure 10 Staticfailure of [0, 9014sspecimen, showingpeel failure within ply 15 (X 1500) agation away from interfaces. As the maximum tensile strain generated within a 90 ~ ply is likely to occur at the boundary closest to the tensile face of the sample, this is not surprising. In some instances it is possible to observe minor cracks opposed to the primary crack propagation direction. If the following crack sequence proposed for plies 13, 14 and 15 is considered, this feature can be explained. A transverse ply crack initiates in ply 15, close to the ply interface nearest the tensile face of the sample. This crack grows rapidly into the ply, leading to stress intensification in the neighbouring plies. It is proposed that at this stage extensive fibre breakage has occurred in both plies 14 and 16, but both plies are capable of bearing high loads. The stress intensification leads to matrix failure in ply 13; again at the interface closest to the tensile face. This in turn causes further intensification of stress in ply 14 and eventual failure of the matrix close to the interface between plies 14 and 15. In this way, ply 15 is subjected to tensile cracking from both ply interfaces. This proposal is supported by the observation that river markings, close to the ply interfaces on either side of the same ply, were frequently oriented in different directions, confirming that transverse ply cracking initiated from a number of different sources. Figure 10 shows cusp formation and this indicated that a significant shear component of stress existed in the plane of the transverse crack, at the crack tip. The orientation of these peel cusps indicates that crack propagation was from left to right. The coarseness of their formation suggests that ply separation was significant, before they were generated. The appearance of peel cusps gives an important clue as to lateral crack growth
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mechanisms within a ply. Frequently a change in cusp orientation is noted. In some instances the point about which this change in orientation takes place can be recognised as the crack origin. Figure 11 shows a distinct change in crack propagation rate occurring in ply 8. In the region marked A, ductile drawing is significant and stripping of matrix material from fibres is clearly evident. In area B the welldefined spherulitic pattern suggests a more rapid crack propagation rate. Figure 12 shows a higher magnification of the ductile failure surface and, significantly, spherulitic features can not be noted even at this very high magnification. The above indicates a sudden slowing of crack propagation rates within the ply. A possible explanation is that transverse cracks are initiated at different positions and on different planes within this 90 ~ ply. The crack propagation directions were laterally opposed. As the crack fronts propagated past one another, a thin strip of the ply was separated from the main fracture surface. The energy of propagation was then transferred to bending of the pulled-out strip and consequently crack propagation changed from brittle matrix cracking to ductile peeling. Extremely rapid crack propagation within plies 8 and 9 is shown in Figure 13. Here, the spherulitic pattern is not visible at x 1000 magnification. However, when this surface was viewed at extremely high magnifications, evidence of the spherulitic formation was again visible. In earlier work 7 it was suggested that crack propagation rates can be categorized as slow ductile, intermediate and brittle, on the basis of spherulitic definition. It is proposed here that, in the case of brittle crack propaga-
Damage development during flexural fatigue: G. Dillon and M. Buggy
Figure 11 Static failure of [0, 9014s specimen, showing change from brittle to ductile failure in ply 8 (x 1500)
Figure 12 Static failure of [0, 9014sspecimen, showing ductile drawing
(x5ooo)
tion, an indication of crack growth rate can be estimated from the magnification necessary to resolve spherulitic features. The spherulitic features are visible because of the differential responses of the constituent crystalline and amorphous regions. Effectively, higher crack propagation rates lead to smoother surfaces requiring higher magnifications to observe this effect. Figure 13 also shows the atypical feature of interlaminar cracking between plies 8 and 9 (x-x). In general, this interface
between two 90 ~ plies remained perfectly intact and plies tended to act as one. The above discussion leads to a fairly clear picture of damage development in cross-plied laminates of statically tested APC-2 laminates. A network of transverse ply cracks builds up in the plies close to the tensile surface, initiation occurring mainly in the regions close to interply layers. Fibre breaks occur on the tensile surface and at ply interfaces, leading to intensification of stresses acting on transverse plies. Progressive fibre failures lead to increased stresses on plies closer to the neutral axis and transverse cracking penetrates deeper into the sample. Compressive cracking occurs in the first few plies and progression into the sample occurs as a straight front. The neutral axis is closer to the compressive surface than in the unidirectional case. This is due to the fact that transverse plies contribute greater support in compression than in tension. Final failure occurs when transverse cracking and fibre failure lead to stress intensity of a sufficient magnitude to fracture the remaining longitudinal fibres. Post-failure delamination and transverse ply pull-out then occur. Finally, it is worth pointing out that crack propagation rates can be accurately traced through each transverse ply. In general the fastest crack propagation rates were noted in the central plies. This was so because they act as a single, much thicker transverse ply, thereby being subject to a lesser degree of crack constraint from 0~ plies. Also, a large degree of damage had already occurred in neighbouring plies before central ply failure. Thus, failure of the central plies is associated with the final rapid failure of the test specimen.
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Figure 13 Static failure of [0, 9014sspecimen, showing interlaminar crack between plies 8 and 9 (x 1000)
Figure 14 Dynamic failure of [0116 specimen tested at 90% of static strength (•
DAMAGE DEVELOPMENT IN FATIGUE FAILURES Fracture surfaces generated by flexural fatigue of unidirectional and cross-plied laminates showed many of the identifiable features found in static failures. However, a number of features particularly associated with fatigue were identified.
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Fatigue failures of unidirectional laminates Figure 14 shows a general view of a sample subjected to a maximum fatigue stress of 90% of its ultimate flexural strength. The sample shown survived 97 load cycles to this stress level. As can be seen by comparison with Figure 1, the fracture surface is similar to that generated in a static test. However, there are two major differences.
Damage development during flexural fatigue: G. Dillon and M. Buggy
Figure 15 Dynamic failure of [0]~6specimen tested at 70% of static strength (x90)
First, the compression failure surface is more complex, particularly near the compression face of the sample. Second, the compression surface is more noticeably curved than in the static case. The complex crack pattern near the surface suggests a gradual build-up of damage by fatigue loading. The horizontal lines x-x, y-y, and z-z are associated with crack arrest points during the last few cycles of loading. The tendency towards compression surface curvature in fatigued laminates is not easily explained but was observed as a general phenomenon. In the static situation the planar character of the compression surface can be explained by the fact that crack propagation occurs in a single stage, and is very rapid. The stored elastic strain energy is released suddenly as unstable cracks initiate from both sample faces. The radius of curvature induced by bending causes cracks on the compression side to propagate at a slight angle to the plane normal to the fibre directions. Once initiated, crack propagation is so rapid that the crack direction does not change until the crack front is very close to the neutral axis. In the dynamic situation crack propagation is much more gradual and directional changes can occur. This may be due to an increasing offaxis stress component, for example shear, as the crack progresses through the compressive region. Supporting this contention is the observation that this effect was more marked in samples subjected to lower fatigue loads. The tensile surface is very similar to that of the statically tested sample. In fact, no significant difference between these two fracture surfaces was noted. However, there was a tendency towards greater fibre pull-out in the fatigued samples. This suggests a more random fibre breakage pattern in the fatigue situation. An angled view of the tensile surface of a sample subjected to a maximum fatigue stress of 70% of the ultimate flexural strength is shown in Figure 15. Again the tensile surface is similar to that of the static case. Single fibres in isolation indicate fibre failure occurring before overall final fracture and these were found to be more in evidence in the fatigued samples. Failure seems to have occurred in intraply bands, suggesting that in this case
crack growth occurred as a number of discrete steps on the tensile side. A number of fibre bundles were found near the tensile face with adhering matrix material, showing ductile shearing. This may indicate either a slowing of crack propagation rates in these regions or, more likely, shear abrasion of matrix material after individual fibre failures had occurred. Clear evidence of abrasion following fibre failure can be seen in Figure 16. Bare fibres were found in a region where crack propagation was rapid. However, as pointed out previously, bare fibres are normally associated with ductile failure. It is considered that a longitudinal crack propagated in a brittle fashion through the area illustrated. Fatigue abrasion of the mating surfaces led to removal of the matrix material surrounding the fibre. As this feature was not noted in statically tested samples, it is suggested that an indication of the amount of longitudinal matrix cracking that has occurred can be estimated from the level of fatigue abrasion observed. As with statically tested samples, the ductility of matrix failure increased towards the neutral axis. As illustrated clearly in Figure 17, fibre damage is much more localized, which may be due to the decreasing radius of curvature. It is also possible that high shear stresses are generated here, as fibre fracture seems to have occurred at 45 ~ to the normal plane in many cases. It is most likely, however, that the fibre stresses tended towards pure flexure as the tensile failures approached the neutral axis. Figure 18 shows a general view of the compression surface of the above sample. This surface can be divided into three distinct regions, in terms of smoothness. Higher magnification views show that the variation in surface smoothness is due to different levels of fatigue abrasion. Figure 19 shows what can be termed low abrasion. Some evidence of the fibrous nature of the material is still visible in this area. In the top left at A, for instance, the general outline of a fibre end can be seen. Figure 20 shows medium level abrasion. Here, the underlying fibres and matrix material have been abraded to such an extent that the individual phases can no longer be distinguished. The entire surface in this region is
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Figure 16 Dynamic failure of [0116specimen tested at 70% of static strength, showing bare fibres indicating fatigue abrasion (x 6000)
Figure 17 Dynamic failure of [0116 specimen tested at 70% of flexural strength, showing fibre breakage near neutral axis (x 1800)
Figure 18 Dynamic failure of [0116 specimen, showing different levels of compressive abrasion (x330)
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Figure 19 Dynamic failure of [0]16 specimen tested at 70% of flexural strength, showing low level abrasion (x 3000)
Figure 20 Dynamic failure of [0116specimen tested at 70% of flexural strength, showing medium level abrasion at interface between plies 2 and 3 (x3300)
covered with damage debris. Figure 21 shows high level abrasion where damage debris has been smoothed off the fracture surface. A series of horizontal lines can also be seen in this fractograph, indicating gradual growth of the compression crack through this region. The dispersion of low, medium and high level abrasion regions throughout the fracture surface was quite complex. In general, a region of medium abrasion is seen close to the compressive surface. This is associated with the relatively rapid initial propagation of a compression crack. The growth rate then slows into the high abrasion region. The smoothness of this region indicates a slow crack propagation rate, suggesting that the greater part of fatigue life is spent in this slow propagation phase. Rapid crack propagation towards the end of fatigue life gives rise to the low level abrasion regions found close to the neutral axis. Finally, a region of pure flexural failure (i.e. no abrasion) occurs at or near the neutral axis. This is associated with the final cycle of testing. It should be noted that regions of high abrasion do not extend across the entire specimen width. Typically, three such regions occur in any one sample, generally
bounded by regions of medium and low abrasion. The obvious implication is that a compression crack is initiated across the sample width and splits into three or four growing cracks, which gradually progress into the sample. Initial cracks grow from the induced fibre damage created by the central loading anvil. In some instances an increase in crack propagation rate occurs in a region of otherwise slow crack growth rate. This is generally accompanied by delamination, as shown in Figure 18 at x-x. Figure 22 shows the shear character of a crack that propagated longitudinally within a ply near the surface. It is thought that growing compression cracks can lead to an increase in matrix shear stresses at the crack tip. The crack grows from the compression surface until encountering an interface. Crack tip shear stresses can then cause the crack to propagate interply. Effectively, an edge notch is introduced to each half of the specimen at the centre and the sample acts similarly to two edge-notch flexure samples. This causes a temporary increase in the rate of crack growth. Previous studies of different crack growth rates in fatigue of composites were not found in the literature. Compressive fatigue fracture has rarely been the subject COMPOSITES Volume 26 Number 5 1995
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Figure 21 Dynamic failure of [0116specimen tested at 70% of flexural strength, showing high level abrasion (x3300)
Figure 22 Dynamic failure of [0116 specimen tested at 70% of flexural strength, showing intraply shear failure near tensile surface (•
of fractographic studies presumably because it has generally been assumed that post-failure abrasion renders fracture surfaces uninterpretable. On the contrary, it is believed that this abrasion provides a useful clue to propagation mechanisms and rates. Interestingly, different abrasion levels were not noted in comparative carbon fibre/epoxy laminates (T300/Cycom 985GT 6095 III CFRP of geometries [0]20, [9015s and [+4515s) studied in the present work. This suggests that progressive damage mechanisms in carbon fibre/epoxy systems are associated with gradual build-up of sub-surface debonds and matrix cracks, whereas in the carbon fibre/thermoplastic system gradual growth of compression cracks represents the main recognisable progressive damage mechanism. Fatigue damage development in unidirectional laminates subjected to low fatigue stresses can be summarized as follows. A compressive crack is induced by load
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member stress intensification. Random fibre breaks occur on the tensile surface. Bands of tensile failure progress into the sample, and a high degree of shear abrasion of failed regions occurs. A number of slowly propagating compression cracks propagate from the original crack. These cracks grow gradually into the sample, giving rise to smooth abraded surfaces. Occasional delaminations occur where high shear stresses are generated at compression crack tips. Final failure occurs when the depth of compression cracking causes fibre stress intensification of a sufficient magnitude to initiate unstable fracture. The degree and type of abrasion can be used to estimate the approximate number of cycles, and consequently the applied stress. Medium to high abrasion is associated with fatigue stresses in the region of 70 to 75% of ultimate strength, whereas low abrasion levels are associated with higher stresses.
Damage development during flexural fatigue: G. Dillon and M. Buggy
Figure 23 Dynamic failure of [0, 9014sspecimen tested at 90% of flexural strength (x43)
Figure 24 Dynamic failure of [0, 9014sspecimen tested at 80% of flexural strength, showing lateral crack growth (x 1000)
Finally, it should be noted that migration of the neutral axis was not observed in unidirectional samples in this study.
Fatigue failures of cross-plied laminates Figure 23 shows a general view of a cross-plied sample subjected to fatigue at 90% of ultimate flexural strength. Comparison with Figure 15 shows that, on a low magnification level, tensile fracture surfaces are similar to those generated in the static situation. In fact, very few differ-
ences between high stress fatigue damaged laminates and statically tested laminates were noted. The degree of delamination in both cases was similar and the same tendency towards crack propagation away from ply boundaries was noted. Figure 24 shows a feature that is believed to be particular to fatigued samples. The feature is a rib, identified 7 as an indication of a change in crack propagation rate. However, though such changes were noted in statically tested samples, ribs of the kind illustrated were not. Closer examination of matrix material
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Figure 25 Dynamic failure of [0, 9014sspecimen tested at 90% of flexural strength, showing striations mirroring underlying fibre surface profile (x20 000)
on either side of the rib shows that crack propagation was in both cases rapid. It is proposed therefore that the rib represents a transverse crack arrest line, between load cycles. The general curvature of this line again suggests that initiation occurred at a ply interface. Features in Figure 25 were initially believed to be fatigue striations. However, further examination showed that these striations simply matched those already present on fibre surfaces caused by the fibre production process. Figure 26 shows a general view of a cross-plied specimen subjected to fatigue at 75% of ultimate flexural strength. As can be seen, the changeover line occurs in the middle of ply 2, suggesting that eventual failure was tensile-dominated. The degree of delamination and fibre pull-out is much greater than in statically tested and high stress fatigue samples. This indicates that gradual buildup of fibre breaks and matrix cracks characterized the
earlier part of the damage process. Towards the end of cycling delamination occurs, leading to stress intensification in surviving fibres. Different crack propagation rates were noted within transverse plies. Figure 27 shows brittle failure 7 in a ply in which extremely rapid crack propagation rates were also noted. Compression cracks generated at the compression sample faces showed some evidence of different levels of abrasion. However, it is believed that surface compression cracks immediately change direction on meeting a ply boundary. A significant amount of delamination in surface plies was noted on the compression side. Also, there was clear evidence of intraply splitting in longitudinal plies. This suggests gradual yielding of matrix material, allowing development of shear cracks within 0 ~ plies. On the tensile side a number of longitudinal matrix cracks were seen to run parallel to the fibres in the 0 ~ plies. Constant-moment flexure tends to induce anticlastic curvature. This curvature is prevented by the presence of the load member in three-point bend flexure, and consequently transverse stresses are induced in the surface plies. These stresses lead to a gradual build-up of tensile surface cracks during fatigue. Evidence was found that suggested that this type of cracking also occurs in longitudinal plies beneath the specimen surface. The appearance of spherulitic features, shown in Figure 28, was particular to fatigued samples. This is very similar to the pattern noted in the interply regions. Both surfaces showed a definite ductile tensile character. The origin of this pattern is not clearly understood. It may be due to gradual progression of a fatigue crack through a spherulite, where fibrils are drawn during successive load cycles. It is also possible that this pattern is associated with a crack propagation rate not previously noted. The fatigue damage development process in crossplied laminates can be summarized as follows. Compression cracking occurs on the sample face in the first few cycles. Sub-surface transverse ply cracking initiates, primarily, near ply interfaces. Fibre breakage occurs in the outermost tensile plies. Intraply longitudinal
Figure 26 Dynamic failure of [0, 9014sspecimen tested at 75% of flexural strength. Angled view at 30~ to show crack route near compressive surface (x47)
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Damage development during flexural fatigue: G. Dillon and M. Buggy CONCLUDING REMARKS
Figure 27 Dynamic failure of [0, 9014s specimen tested at 75% of flexural strength, showing intermediate brittle failure on pulled-out 90~ ply
(x4800)
matrix cracks are initiated by high interfacial shear stresses parallel to the ply interfaces. Longitudinal matrix cracks are also initiated by high transverse stresses generated in the 0 ~ plies. Growing compression cracks and transverse ply cracks penetrate interply regions giving rise to delaminations. All the above mechanisms continue to operate until unstable cracking occurs in the final cycle.
The present study illustrates the usefulness of scanning electron microscopy in identifying failure mechanisms in composite laminates. Previous investigators have shown that a good deal of useful information on static failure mechanisms can be attained using this technique. This work shows that many of the features identified in the static situation can also be seen in samples subjected to flexural fatigue loading and that additional features particular to the dynamic loading situation can be recognised. Fatigue striations were not seen in the form normally associated with polymeric materials, except on a coarse scale at very high stress levels, and then only in unidirectional samples. More useful information on the fatigue history of a laminate can be gained by studying the level of fatigue abrasion in compression cracks generated directly beneath the load member. Higher levels of abrasion indicate longer lives and therefore lower stress levels, in the region of 70% of the ultimate strength of the laminate. An indication of the level of longitudinal ply matrix cracking occurring can be estimated from the degree of matrix fatigue abrasion noted. Cross-plied laminates showed extensive transverse and longitudinal ply cracking and interlaminar shearing. The usefulness of the technique is somewhat limited by laminate geometry, samples containing angled plies showing extremely complex shear failures, although shearinitiated progressive compression failure was identified. Nonetheless, it is felt that more attention should be focused on scanning electron microscopy as a means of identifying failure mechanisms in composite materials.
Figure 28 Dynamic failure of [0, 9014s specimen tested at 75% of flexural strength, showing coarse spherulitic fracture on 90~ ply (x4800)
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Damage development during flexural fatigue: G. Dillon and M. Buggy REFERENCES 1 Talreja, R. 'Fatigue of Composite Materials', Technomic Publishing Co. Inc., Westport, CT, 1987 2 Hartness, J.T. in 'Proc. 28th Nat. SAMPE Syrup.' 12-14 April 1983, p. 535 3 Kim, R.Y. and Hartness, J.T. in 'Proc. 29th Nat, SAMPE Syrup.' 3-5 April 1984, p. 765
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Buggy, M. and Dillon, G. Composites 1991, 22 (3), 191 Purslow, D. Composites 1981, 12 (4), 241 Purslow, D. Composites 1986, 17 (4), 289 Purslow, D. Composites 1987, 18 (6), 365 Purslow, D. Composites 1988, 19 (2), 115 Purlsow, D. Composites 1988, 19 (5), 358 Croman, R.M. in 'Proc. ICCM and ECCM', London, UK, July 1987, p. 4.76 11 Bader, M.G. and Johnson, M. Composites 1974, 5, 58 4 5 6 7 8 9 10