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epoxy laminates subjected to thermal cycling
COMPOSITES SCIENCE AND TECHNOLOGY Composites Science and Technology 64 (2004) 1725–1735 www.elsevier.com/locate/compscitech
Influence of oxidative environments on damage in c/epoxy laminates subjected to thermal cycling M.C. Lafarie-Frenot *, S. Rouquie Laboratoire de Mecanique et de Physique, des Materiaux, – UMR CNRS No. 6617, ENSMA, 1 Av. Clement Ader – B.P. 40109, 86961 Futuroscope-Chasseneuil Cedex, France Received 3 March 2003; received in revised form 12 January 2004; accepted 15 January 2004 Available online 20 February 2004
Abstract The aim of this study is to characterise damage processes in carbon/epoxy laminates submitted to thermal cycling in neutral (nitrogen) and oxidative (air, oxygen) atmospheres. Observations of the polished edges of the specimens by optical microscopy and SEM have revealed the presence of matrix shrinkage and fibre/matrix debondings only on samples tested under oxidative atmospheres. Quantitative analyses of transverse matrix cracking by means of microscopy and X radiography have shown the accelerating effect of an oxidative atmosphere. Moreover, the higher the oxygen concentration, the more significant the acceleration of the cracking development. These results are analysed in terms of matrix oxidation, occurring at the highest temperatures of the cycle, and fatigue matrix cracking, occurring at the lowest ones and due to the prevented differential expansions of the plies. During a thermal cycling test performed in an oxidative atmosphere, a coupling effect between these two damage mechanisms would result in the acceleration of the micro cracking observed. 2004 Elsevier Ltd. All rights reserved. Keywords: B: Durability; B. Matrix cracking; B. Thermomechanical properties; Environmental degradation; B. Fatigue
1. Introduction The present study was conducted in the framework of the French ‘‘Aeronautical Supersonic Research’’ program. The objective of this program is to direct the fundamental research in fields where scientific and technological projections are essential for the design of a new generation of supersonic aircraft. Compared to Concorde, the weight saving required on the aircraft structure is estimated at about 30%. This requirement makes necessary the use of organic composite materials for the manufacturing of most of the structure (Concorde was primarily made of aluminium alloys). This new future supersonic aircraft, designed to fly at Mach 2, will have a lifetime from 3 to 4 times longer than that of the Concorde, corresponding to a minimum of 80,000 h, i.e. approximately 20,000 flights. In this type of *
Corresponding author. Tel.: +33-5-4949-8229; fax: +33-5-49498238. E-mail address: [email protected] (M.C. Lafarie-Frenot). 0266-3538/$ - see front matter 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.compscitech.2004.01.005
application, the materials constitutive of the aircraft structural parts can be subjected to cyclic mechanical loadings and to temperature variations of great amplitude, depending on the phase of the flight, subsonic or supersonic. The objective of a significant reduction in mass, compared to Concorde, would lead to a massive use of composite materials, and particularly of carbon fibre – polymeric matrix laminates. In this context, a better prediction of the long-term durability of composite materials with organic matrix is essential. In particular the study of the coupled effects of mechanical cyclic loads, of temperature variations and of a more or less oxidative environment on the damage of these materials remains to be made. Within the framework of the American program ‘‘High Civil Speed Transport’’, project at Mach 2.4, Arendt et al. [1] described a campaign of tests meant to estimate the durability of some composite materials, candidates to this type of application. Averaged cyclic thermo mechanical conditions are simulated, consisting of the superposition of temperature cycles ()36, 200 C)
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and of stress cycles, the period of which is the mean duration of a flight. This approach leads to very long tests, is very expensive, and requires making a priori the choice of definite materials. On the other hand, Gates [2] proposed another procedure, consisting in the definition of accelerated tests by increasing some critical values of parameters characteristic of physical ageing and chemical degradation of the materials considered. He insists on the need of understanding and characterising the coupling that can appear between these parameters, under the combined effects of a mechanical loading, of the temperature and of the environment. When composite laminates with long continuous fibres are subjected to temperature variations, the mismatch of thermal expansion coefficients of fibres and matrix as well as the difference of ply orientation in the lay up, are such that local stresses appear, which can take part in the degradation of the laminate. When these thermal variations are cyclic, they induce, at the ply level, cyclic stress variations which can be compared, at this scale, to a fatigue phenomenon. Various types of damage similar to those observed in mechanical fatigue result from these cyclic stresses, like transverse matrix cracking, fibre/matrix debonding and delamination [3,4]. When these temperature variations occur in the presence of an oxidative environment, other damaging phenomena appear due to matrix oxidation. Concerning thermoset epoxy polymers, the oxidation phenomenon that occurs during high temperature exposure in oxidative atmospheres, was studied through isothermal ageing experiments by Madhukar et al. [5] and Colin et al. [6] The oxidation of the epoxy matrix results in a loss of mass and a reduction in volume of the matrix, inducing shrinkage of the matrix with respect to the fibres [7]. The study carried out by Colin et al. [8], on epoxy and bismaleimide resins, emphasises for both materials the superficial character of oxidation, which results in the creation of a low thickness oxidised layer on the free edges of the samples subjected to ageing. During ageing, the thickness of this layer grows quickly towards an asymptotic value depending on the studied material and on the temperature of ageing. The detection and the characterisation of this oxidation were also completed by micro hardness Vickers tests [9] on various composite laminates with carbon fibres and different thermoset or thermoplastic matrices. These authors have shown that the oxidised layer has a smaller Vickers hardness compared with that of the virgin material and that this reduction in hardness depends on the atmosphere oxygen concentration. Few works deal with the long-term behaviour of composites subjected to high temperature under oxidative atmosphere. Bowles [10] studied the ageing of T650-35/ PMR15 composites at various temperatures and experimentally found a linear relation between the properties in
compression of the laminates and the thickness of the oxidised layer. The prediction of matrix cracking growth under thermal cyclic variations was not studied as much, as its mechanical fatigue loading counterpart. Favre et al. [11] studied the damage behaviour of various composite materials, with carbon fibres and thermoset matrices (epoxy, cyanate or bismaleimide), subjected to thermal cycling. These authors highlighted an interaction between the matrix micro-cracking of the laminate layers due to the cyclic thermal stresses and oxidation already observed under isothermal ageing. However, as these authors noted, the knowledge of the mechanisms of both crack accumulation and of oxidation is necessary before a relevant prediction can be made. The aim of this experimental study is to characterise the damage growth in carbon/epoxy laminates subjected to cyclic variations of temperature ()50 C/180 C) and to specify the influence of the considered atmosphere: neutral (nitrogen) or oxidative (dry air or pure oxygen). The longer-term goal is to determine protocols of ageing tests, accelerated compared to the tests of thermal cycling presented in this paper. This ambitious objective requires a good preliminary understanding of the oxidation and damage mechanisms, their respective ranges of validity and their interactions.
2. Experimental conditions 2.1. Materials and preparation Thermal cycling tests have been performed on continuous carbon fibre and epoxy matrix composites, selected by EADS in a previous study. The specimens are parallelepipedic, of dimensions 35 · 25 mm2 , the length direction being referred to as the 0 axis (cf. Fig. 1). Four stacking sequences have been investigated: unidirectional [0]8 (UD), cross ply laminates [03 /903 ]S , [453 /)453 ]S , and [45/0/)45/90]S (quasi-isotropic – QI) (their respective thicknesses are 1.04, 1.68, 1.68, and 1.16 mm). The cycle of polymerisation has been optimised by EADS CCR in order to obtain a stable material, to avoid any effect of post cure and change of properties during thermal cycling (cf. Fig. 2).
Fig. 1. Specimen geometry and observations of the polished edges.
M.C. Lafarie-Frenot, S. Rouquie / Composites Science and Technology 64 (2004) 1725–1735
Fig. 2. Polymerisation cycle and post cure of studied composite.
2.2. Methods First, the samples have been smoothly polished on two perpendicular sides. They were examined prior to cycling in order to ensure that there were no initial cracks or defects on the free surfaces. 500 thermal cycles have been prescribed and regularly the samples are removed from the thermal cycling oven to observe damage development: on the free edges of the specimens by optical microscopy and scanning electron microscopy (SEM) and inside by means of penetrant enhanced X radiography. As a precaution the samples were not reused after SEM observations or X radiographs. In the first case, the polished edges were coated with a thin film of gold to avoid any charging effect during SEM observations and, in the second one, zinc iodide solution was used to enhance the contrast of the cracks. In both cases, physical or chemical modifications could induce changes in further damage development. 2.3. Experimental device For this study a specific thermal cycling device has been developed in order to control both the environment and the temperature of the test (Fig. 3).
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Polished and dried samples are placed vertically on racks (cf. Fig. 3(a)), in order to have homogeneous gas flow and temperature distribution around the specimens. Then, they are arranged into a specific enclosure (cf. Fig. 3(b)) having low volume and thermal inertia, allowing constant gas flow with an overpressure of 20 mbar. Various gas mixtures, of controlled relative contents, are obtained from industrial dry gases and pumped into the enclosure. Finally, this device is placed inside a thermal equipment allowing controlled temperature variations to be prescribed (cf. Fig. 3(c)). To assess the influence of oxygen on material ageing, the tests presented here were conducted in three atmospheres: neutral (pure dry nitrogen) and oxidative (dry air reconstituted by 22% O2 + 78% N2 , and pure dry oxygen). The prescribed thermal cycles were triangular, the maximum and minimum temperatures being, respectively, of 180 and )50 C with constant cooling and heating rates of 4 C/min. Fig. 4 shows an example of the real testing thermal cycle prescribed, temperature being measured with a thermocouple embedded in the specimen. It can be noted that the glass transition temperature of this composite material, measured by DSC, has been found around 200–210 C, slightly varying with the material batch considered. In order to accelerate the damage and ageing processes, the maximum temperature of the cycle (180 C) has been chosen higher than the real temperature of the material during a supersonic flight (100 C/130 C) but smaller than Tg. 2.4. Thermo elastic calculations of ply stresses Thermo-elastic calculations [12] give an idea of the level of thermal stresses present in each ply of the laminate for a given temperature. These stresses stem from the mismatch of thermal expansion coefficients of plies with different orientations. The lay up restrains differential thermal deformation of the plies and induces
Fig. 3. Environmentally controlled thermal cycling device.
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Fig. 4. Experimental record of a temperature cycle ()50 C/180 C).
period is equivalent to a sort of thermally driven ‘‘fatigue’’ with DT analogous to the change in stress as previously noted by Hancox [13]. It is important to point out that in an unidirectional laminate, made up of eight plies of same orientation, the thermal stresses are only due to the difference of thermal expansion coefficients between fibre and matrix and not, as for the other laminates of this study, between plies of various orientations.
3. Experimental results 3.1. Damage mechanisms due to thermal cycling
Fig. 5. Thermal cycling and corresponding thermal transverse ply-stresses.
thermal stresses in each oriented layer. Moreover, the magnitude of thermal stresses in the material increases as the temperature of the material deviates from the stress-free temperature. For example, in the case of the thermal cycling between )50 and 180 C, in the different layers of the cross ply laminate [03 /903 ]S , the transverse thermal stresses lie between 0 and 45 MPa. In Fig. 5, the transverse thermal stresses (r22th ) are put in correspondence with the variation of temperature: the cyclic temperature variation induces cyclic transverse stress variations, in each ply of the laminate. This cyclic ‘‘thermal loading’’ with a 2 h
Thermal cycling tests induce various types of damages, depending to a certain extent on the orientation and the thickness of the layers in the lay-up, but mainly on the environment of the specimen. In order to compare damages according to the location of the plies in the stacking sequence of the laminate, microscopic observations have been conducted on two perpendicular sides of the samples. Throughout thermal cycling tests, the observations of the free edges have shown three types of damages: permanent deformation of the matrix due to its shrinkage (Fig. 6(a)), debonding between fibres and matrix (Fig. 6(a)), and matrix cracking (Fig. 6(b)). Considering matrix cracking, two kinds of cracks can be observed. ‘‘Transverse cracks’’ are visible in the plies where the fibre direction is not parallel to the observation plane and, in this study, we call ‘‘transverse cracks’’ the only cracks which cross the whole thickness of the layer concerned (Fig. 6(b)). These cracks are propagating through the matrix but with a direction guided by the fibres; for example, in 90 plies sur-
Fig. 6. Different types of damages. [03 /903 ]S oxygen, 100 cycles.
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rounded by 0 plies, they appear globally perpendicular to the layer interfaces (Fig. 6(b)). Moreover, especially in the thick central layer of the [03 /903 ]S laminates, other kinds of cracks may exist; we called them ‘‘small cracks’’ (Fig. 6(b)). The length of these small cracks can vary between a few fibre diameters to half the thickness of the layer and appear more or less slanting on microscopic observations (Fig. 6(b)). Depending on the experimental conditions (stacking sequence, environment, etc.) and the number of cycles, either none, or some, or all these types of damages could be observed. The damage features as well as their variations, according to the different studied parameters, are described in the following sections. 3.1.1. Matrix shrinkage The example of the [03 /903 ]S cross ply laminate is shown in Fig. 7. In this figure, SEM pictures of the free edges of specimens that have sustained 500 thermal cycles, respectively, in nitrogen, in air and in oxygen, are compared. Note that these pictures have been taken with a 45 tilt of the samples. Comparing the three pictures in Fig. 7, one can see that a significant shrinkage of the matrix only exists on the edges of the specimens tested in oxidative environments. This matrix shrinkage is visualised as dark and deep hollows, imprinted on the initial plane of the polished edge. Moreover, the wider the matrix areas concerned, the deeper are the hollows (Fig. 7). In nitrogen, the very small differences in level observed through the scanning electron microscope between fibres and matrix are similar to those observed on every polished specimen edge before any test. Therefore, the aspect of the edge visualised in Fig. 7(a)) is associated with the polishing stage, whereas the matrix shrinkage shown in Figs. 7(b) and (c) is due to an oxidation process. The matrix shrinkage is observed on the sample surfaces for each laminate tested in oxidative environment (air or oxygen). Whatever the stacking sequence is, these shrinkages appear distributed on each of the free edges. However, deep hollows are systematically observed in rich matrix areas that are present at interfaces between layers of different orientations. The unidirectional case is special because all layers have the same orientation, annihilating the stresses due
Fig. 7. SEM pictures (45 tilt) of the polished edges of [03 /903 ]S laminates subjected to 500 thermal cycles: (a) nitrogen, (b) air and (c) oxygen.
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to the mismatch of thermal expansion between layers of different orientation. However, on such laminates, the matrix shrinkage is also present when the test has been performed in oxidative atmosphere. These residual deformations of the matrix have the same characteristics as the shrinkage observed in the other laminates considered (Fig. 8). 3.1.2. Debonding and crack initiation When matrix shrinkage occurs (i.e. in oxidative environment), the high strain gradients present in matrix areas, close to fibres which have a very high stiffness, lead to high local stresses that can be favourable to debonding between fibres and matrix or to crack initiation as shown in Figs. 9 and 10.
Fig. 8. MEB observation (45 tilt) of the edge of a [0]8 sample after 500 thermal cycles in oxygen.
Fig. 9. QI laminate at 500 cycles in oxygen (45 tilt).
Fig. 10. SEM observations of a [03 /903 ]S specimen subjected to 100 thermal cycles in air.
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As an example, Fig. 9 shows a SEM picture of the edge of the quasi-isotropic laminate [45/0/)45/90]S subjected to 500 thermal cycles in oxygen. In that figure, one can see a deep hole in the matrix, which is located at the interface between the central 90 layer and a 45 ply. In that case, the matrix contraction has induced numerous fibre/matrix debondings in the circumference of the hole. For each test, a first interruption has been made after 100 thermal cycles only. When the tests have been performed in oxidative environment, it appeared that, after 100 cycles, matrix shrinkage is already well developed. Consequently, fibre/matrix debonding, which seems to be connected with this type of damage, has been observed very early during the test. For example, in Fig. 10(a), some small fibre/matrix debondings are present in a [03 /903 ]S specimen subjected to only 100 thermal cycles in air. Moreover, observations of the same specimen, at the same stage of the test, have shown short matrix cracks (Fig. 10(b)). However, one cannot yet be sure that these cracks are initiated from fibre/ matrix debonding, propagating from one such debonding to another? – or if they are induced by another matrix fracture mechanism. 3.1.3. Transverse matrix cracks In Fig. 11, the SEM aspects of transverse cracks present in two [03 /903 ]S specimens after 500 thermal cycles, either in oxygen (Fig. 11(a)) or in nitrogen (Fig. 11(b)), are compared. Once more, important differences in relief of the specimen edges appear: they remain very flat in specimen tested in nitrogen, whereas they are quite uneven in specimen tested in oxygen. These pictures show that the crack opening is larger in O2 than in N2 . In oxygen, the cracks are wide, well opened, allowing observations of the cracked surface of the material: in that case, the fracture pattern appears very brittle, most of the fibres are naked with no traces of residual matrix on them (Fig. 11(a)). Comparatively the cracks propagated in nitrogen are more superficial and narrow (Fig. 11(b)). In that case, the crack opening is so small that the fracture pattern cannot be observed
Fig. 11. SEM observations of transverse ply crack (45 tilt) in [03 /903 ]S subjected to 500 thermal cycles: (a) oxygen and (b) nitrogen.
by SEM and therefore cannot be described from such observations. Equivalent observations of transverse cracks obtained by thermal cycling in air have shown intermediate trends, between those of specimens tested in oxygen and in nitrogen. However, the transverse cracking features (large crack opening, brittle aspect of the fracture, etc.) have been found closer to those of the cracks developed in oxygen: that observation emphasises the influence of the oxidative environment on the transverse cracking mechanisms. To conclude this section concerning microscopic observations of damage due to thermal cycling in neutral or in oxidative environment, the main results can be summarised as following: • Matrix shrinkage is observed on surfaces of specimens tested under oxidative atmosphere only. • In that case, debonding between fibres and matrix are more numerous (and that would be probably related to the above point). • Compared to nitrogen, in oxidative environment (air or oxygen), after a given number of thermal cycles, transverse matrix cracks are more numerous, wider opened, and are increasing in number much faster with numbers of thermal cycles. 3.2. Quantitative analysis of cracking As was shown in the above sections, the damage features as well as the kinetics of its development with number of cycles depend on the specimen environment. In a first step and in order to evaluate more quantitatively these differences, we have focused our study on the sole transverse matrix cracking. Moreover, in this paper, we have chosen to present only the results concerning the cross-ply laminate [03 /903 ]S . Indeed, as well known in fatigue, such a lay-up, with ‘‘thick’’ layers of 3 or 6 plies, is favouring transverse cracking, compared to stacking sequences with thinner layers. The crack measurements (number, length, etc.) are therefore easier and more precise in that laminate. However, in the other stacking sequences considered in this study, the quantitative characterisation of the damage has shown tendencies similar to those emphasised in the [03 /903 ]S laminate and described in the following sections. 3.2.1. X-ray observations The microscopic observations of the polished edges of the samples are complemented by X radiographs which reveal the crack distribution inside the specimen (cf. Figs. 12 and 13). The use of a zinc iodide solution is required to enhance the picture contrast. For crack length measurement, the resolution of this method has been estimated at 1 mm. We have noticed previously that, in the [03 /903 ]S laminate, two types of matrix cracks coexist: ‘‘transverse
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Fig. 12. X radiographs of [03 /903 ]S at 100, 300 and 500 cycles in air.
Fig. 13. Comparison of X radiographs at 500 cycles in nitrogen (a), air (b) and oxygen (c).
cracks’’ which span the entire ply thickness and are wide open, and oblique ‘‘small cracks’’. Although we observed numerous small cracks in the inner layer of that laminate by optical microscopy, none appears on the Xray pictures. This is probably due to their short length, which could be less than the resolution of the X radiographs. In air (cf. Fig. 12), transverse cracking develops significantly between the different observation steps. At 100 cycles, only few short cracks are counted in the central layer. Then, at 300 cycles, these cracks have propagated and increased in number, and transverse cracks have appeared in the external layer, spanning the whole sample length. Finally, at 500 cycles, the transverse crack distribution is quasi-periodic and probably saturated in number. In the central layer some cracks occupy half the width (12.5 mm) but for the most part their length lies between 2 and 10 mm. In the external layers the cracks are more regularly distributed through the sample which appears evenly damaged along its whole length (cf. Fig. 12). In Fig. 13, the X-ray pictures obtained after 500 thermal cycles, in the three studied environments are compared. In nitrogen (Fig. 13(a)), the damage built up at 500 cycles is quite small. Few short cracks (whose length is less than the quarter width of the layer) are seen in the central layer whereas no crack is revealed in the external layers by the microscopic observations. In oxygen the transverse cracking development is similar to that observed in air, but with faster kinetics. In Fig. 13(c) obtained after 500 thermal cycles in oxygen, one can observe that the sample is uniformly damaged, the crack distribution appearing very regular, and more regular than in air (Fig. 13(b)). Moreover, at that step of the thermal cycling, some small grey areas can be seen on X-ray pictures, surrounding some matrix cracks and corresponding to
partial delamination between central and external layers. The data concerning the transverse crack numbers obtained either from the microscopic observation of the polished edges of the samples or from X radiographs have been compared. It appeared that both data are very close when the ‘‘transverse cracks’’ only are considered on microscopic observations (neglecting the ‘‘small cracks’’) and when only the cracks very close to the edge on X radiographs were counted. As a consequence, we have chosen to characterize the damage level by two measurements: the ‘‘transverse crack’’ density measured on the edge of the sample and the total cracked surface measured on the X-ray pictures. 3.2.2. Transverse crack density In Fig. 14, the transverse crack densities measured on the polished edges of the [03 /903 ]S samples cycled in nitrogen, air and oxygen are plotted versus the number of cycles. They are expressed in number of cracks per centimetre, and distinguished according to the position of the considered layer in the stacking sequence: either in the six ply central layer, or in the three ply external layer. In Fig. 14, the full lines correspond to the average development of the crack densities in the central layer, and the dotted lines to that in the external layers. The crack density measurements have been repeated twice, on two different specimens, and for the external layer, each value of crack density plotted in this figure is the average of measurements made on both external layers. It can be seen in Fig. 14 that the scattering on the measurements is rather small, and low enough to consider the average curves as representative of the crack density growth throughout the thermal cycling test. Comparing the three environmental conditions, one can see in Fig. 14 that the transverse cracking development is much faster and more important in an oxidative
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10
oxygen
N cracks/cm
8
air 6
4
2
nit r ogen 0 0
50
100
150
200
250
300
350
400
450
500
Number of cycles Fig. 14. Edge transverse crack densities vs number of thermal cycles ()50, 180 C), [03 /903 ]S , in nitrogen (dark grey), air (thin black line) and oxygen (large black line).
atmosphere (oxygen and air) than in a neutral one (nitrogen). In nitrogen, cracks have been observed in the central layer only, by both optical microscopy and X radiography. There was no crack in the external layers, even after 500 thermal cycles. In that environment, the first cracks appeared between 200 and 300 cycles and they increased in number slightly, up to a level of about 2 cracks/cm at the end of the thermal cycling test. In air, the first cracking onset occurred much earlier than in nitrogen: after 100 thermal cycles, the crack density measured in the central layer is already equal to 3.3 ± 1 cracks/cm. In that layer, the increase in the crack number is fast, up to 300 cycles when saturation is reached with a value of 7.4 ± 0.8 cracks/cm. In the external layers, the cracks initiated later, between 100 and 200 cycles; and their number regularly increased up to a level of 6.9 ± 0.9 cracks/cm at 500 cycles. Considering the data scatter, it appears that the damage level reached at 500 cycles in air is quasi-identical for both central layer and external layers. Nevertheless, two kinds of crack development seem to be involved: a saturation state is reached in the central layer whereas, in the external layers, the increase in the crack number is still fast at the end of the test. We could assume that for a number of cycle higher than 500, the crack density in the external layers is higher than that in the central layer. In oxygen, the kinetics of transverse cracking due to thermal cycling is the fastest we have observed. After only 100 cycles, the crack density in the central layer is 6.1 ± 0.1 cracks/cm, and quasi-saturated, with a saturation value of 7.7 ± 0.6 cracks/cm at 300 cycles. This saturation value can be considered as identical to the one measured in air, in the same central layer. If we examine the cracking development in external layers, we
can assume that the cracks initiated before 100 cycles, since the crack density is equal to 1.6 cracks/cm as early as 100 cycles; then, the number of cracks increased very quickly, up to a crack density of 9.6 ± 0.2 cracks/cm at 500 cycles. The comparison of the transverse crack growth according to the environment of the thermal cycling test emphasises the accelerating effect of an oxidative atmosphere. Moreover, the higher the oxygen concentration, the more significant the acceleration of the damage processes. 3.2.3. Cracked area In order to represent the damage level of each sample, the total cracked areas have been evaluated from X radiograph enlargements. For each crack visualised on an X-ray picture, a cracked area value has been obtained by multiplying its length by the thickness of the layer concerned. The total cracked area is then obtained by summation over all transverse cracks present in each layer of the sample. The cracked area values presented in Table 1 represent the damage levels of [03 /903 ]S specimens (35 · 25 · 1.68 mm3 ), at the end of each test, after 500 thermal cycles. The comparison of the cracked area values according to the test environment points out, very clearly, that the damage level caused by matrix transverse cracking is very low in nitrogen. The trend seen above by measuring Table 1 Cracked area (mm2 ) measured in [03 /903 ]S samples (35 · 25 · 1.68 mm2 ), subjected to 500 thermal cycles ()50 C/180 C)
[03 /903 ]S
Nitrogen
Air
Oxygen
28.5 ± 1
500 ± 18
580 ± 16
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the edge crack densities is enhanced: the cracked surface is higher in oxygen (580 mm2 ) than in air (500 mm2 ) and much larger than in nitrogen (28.5 mm2 ). The relatively small difference of cracked surface values between air and oxygen is due to differences in crack length in the central layer and in crack density in the external layers. Indeed, we have noted in the preceding section that, at 500 cycles, the central layer is saturated with transverse cracks in both oxidative environments: the corresponding cracked area values therefore depend only on the crack lengths, which are smaller in air than in oxygen. Moreover, in the external layers, almost all cracks span the length of the sample, but, as seen in Fig. 14, the crack density at 500 cycles is lower in air than in oxygen. The thermal cycling tests performed on other laminates highlighted the same trends, whatever the stacking sequence. The damage processes are always highly accelerated by an oxidative environment and even more under oxygen than under air. The differences essentially consist in the respective values of crack density and length, depending on the layer thickness, its location in the stacking sequence, etc. For example, we found that the thinner the layer is, the later the first crack appears and for a given cycle number, the shorter the cracks are. These observations are similar to those noted in fatigue of cross-ply laminates with different 90 ply thicknesses [14]. However, in oxidative environments, we found a faster crack development in external layers than that expected by considering their thickness and their location: that could highlight a local influence of the oxidation process. These remarks are similar to those made by Herakovich et al. [15], who studied different graphite/ polyimide lay-ups subjected to various kinds of thermal exposures, and who found that the density of micro cracks depends on fibre orientations, stacking sequence, edge effects and cooling rate. A 2D thermo-elastic calculation of stresses which develop in the layers during the thermal cycling shows that they are identical for the three laminates [03 /903 ]S , [453 /)453 ]S at [45/0/)45/90]S (QI). However, we have noted that, in oxidative environment, the cracking damage induced by 500 thermal cycles are very dependent on the lay-up: in particular, the cracked surface area measured in the [03 /903 ]S laminate (e.g. in oxygen: 580 mm2 ) is twice higher than in the [453 /)453 ]S laminate (210 mm2 ) and much higher than in the QI sample (6.5 mm2 ). The stress calculations take into account only the number of plies of same orientation but neither the thickness of the layers, nor their constraining and nearness to the oxidative environment. Moreover, because of the viscous behaviour of the polymer matrix which is enhanced at high temperature, stress relaxation might be significant over the times involved in these tests. Consequently, the thermo-elastic calculations which only give the level of the initial mean stresses in the plies are inadequate to analyse and discuss the
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experimental results that might depend on the local state of stresses on the free edges of the specimen. Because the oxidation process is related to the diffusion rate of oxygen in the matrix, a local approach might be more efficient to enhance the understanding of the role of the lay-up on the damage development in oxidative atmosphere.
4. Discussion The optical and SEM microscopic observations of the surfaces of samples subjected to thermal cycling ()50 C/180 C), have shown a significant effect of oxygen on the damage mechanisms: it results in a matrix shrinkage, seen on all the samples tested in air and oxygen whereas the surfaces of the specimens tested in nitrogen remain as flat as they were after polishing [16]. The experimental protocol is identical in oxygen, air and nitrogen: same preparation, same time of preliminary drying, same polishing and same thermal cycling. The only parameter changing between the series of tests is the oxygen concentration of the test environment. The differences described above can therefore be associated to this parameter. In oxidative atmospheres, the shrinkage of the matrix and the debonding between fibres and matrix could be induced by the oxidation of the matrix, very active between 120 and 180 C (cf. Fig. 15). Indeed, the superficial oxidation of the epoxy resin induces a reduction in volume [8] resulting in a shrinking of the matrix on the composite surfaces. The significant deformations associated with that shrinkage could induce high stresses on the fibre/matrix interfaces; these stresses can lead to interfacial fracture. Moreover, the oxidised matrix has probably degraded strength properties as well as a lower tenacity. A recent study on oxidised thin films of the same epoxy resin has shown that the failure stress and strain values decrease as the ageing duration increases [17]. As seen on Fig. 10, the zones of matrix shrinkage, characteristic of oxidation phenomena, seem to be favourable sites of debonding between fibres and matrix, and onset of matrix cracking. Temperatur Temperature
oxidation ox 120˚C fa t igue fatigue Time cr acking cracking
Fig. 15. Thermal cycling in oxidative atmosphere and phenomena involved.
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This primer damage mechanism due to oxidation would lead to an easier and earlier crack initiation observed in Fig. 14 when comparing oxygen, air and nitrogen transverse cracking history. The onset and propagation of transverse matrix cracks probably occur at the lowest temperatures of the thermal cycle, where the stresses due to the prevented differential expansions of the plies are most significant. Timmerman et al. [18] have studied various types of carbon/epoxy cross ply laminates [03 /903 ]S performing thermal cycling down to cryogenic temperature. They have observed that, as the temperature decreases, the increase of residual stresses due to the difference of coefficients of thermal expansion between fibres and matrix and between plies of various orientations, are inducing micro cracks in the composite plies. The same remark was made by Pederson and Gillespie [19], who studied isothermal ageing at 177 C of BMI and polyimide unidirectional and cross ply laminates. They show that due to thermal expansion mismatch of the plies, residual ply-stresses present in the cross ply laminate, but absent of the unidirectional specimen, provide the driving force for crack propagation. In neutral environment such as nitrogen, the poor level of cracking observed at the end of the test could only be due to the cyclic variations of ply-stresses induced by the thermal cycling. It could be equivalent to that induced by 500 cycles of ‘‘fatigue’’, with the same ply-stress amplitude. In oxidative environment, this crack growth at the lowest temperature of the thermal cycle could be accelerated by the higher brittleness of the matrix and the interfaces due to oxidation. Moreover, with each onset of a new crack, reactive surface increases, which accelerates the process of local oxidation and damage [20,21]. In that case, a coupling effect between oxidation and matrix cracking would result in an acceleration of the damaging processes.
5. Conclusion This study concerning the damage of carbon/epoxy laminates subjected to cyclic variations of temperature ()50 C/180 C), highlights the influence of the test atmosphere: neutral (nitrogen) or oxidative (air or oxygen) on the damage mechanisms. Observations have shown: • Residual deformations of the matrix, seen as hollows on the sample surfaces and seen only in oxidative atmosphere (air and oxygen). This matrix shrinkage was observed in every lay-up, within the layers as well as at the interface between layers of different orientations. • More numerous and deeper debondings between fibres and matrix in air and oxygen compared to nitrogen.
These two damage mechanisms of the specimen surface could be induced by the matrix oxidation which is very active above 120 C in oxidative atmospheres. These zones, weaken by the oxidation, are favourable sites for crack onset. The observations made during thermal cycling tests have also shown matrix cracking much more developed in oxidative environment than in nitrogen. The coupling between oxidation, due to oxidative atmosphere, and thermo-mechanical cyclic stresses, due to temperature variations, could accelerate the damage process, all the more than the oxygen concentration is high: in that study, the matrix cracking has been found more developed in oxygen than in air.
Acknowledgements The authors acknowledge the financial support of the French Research Department, the French Transport Department, and EADS (CCR Suresnes, France) for supplying the composite plates and samples.
References [1] Arendt C, Brunner M, Wang J, Cebeci T. Long term durability testing of polymer composite materials. In: Massard T, Vautrin A, editors. 12th International Conference on Composite Materials (ICCM12), Paris, France, 5–9 July; 1999. [2] Gates TS. Durability assessment of polymeric composites for high speed civil transport. In: Proceedings of the DURACOSYS 99, Brussels, Belgium, July 11–14; 1999. [3] Henaff-Gardin C, Lafarie-Frenot MC, Desmeuzes JL. Modelling of crack evolution under cyclic thermal loading in cross-ply laminates. In: Poursartip A, Street K, editors. 10th International Conference on Composite Materials, Whistler, Canada, August 14–18, Woodhead Publishing Limited; 1995. p. 447–54. [4] Henaff-Gardin C, Goupillaud I, Lafarie-Frenot MC. Evolution of matrix cracking in cross-ply CFRP laminates: differences between mechanical and thermal loadings. In: Cardon AH, Fukuda H, Reifsnider KL, Verchery G, editors. DURACOSYS 99, Brussels, Belgium, July 11–14, 1999. Recent developments in durability analysis of composite systems, Balkema; 1999. p. 69–76. [5] Madhukar MS, Bowles K, Papadopoulos DS. Thermo-oxidative stability and fiber surface modification effects on the in plane shear properties of graphite/PMR 15 composites. J Compos Mater 1997;31(6):596–618. [6] Colin X, Marais C, Favre JP. Damage/weight loss relationship of polymer matrix composites under thermal ageing. In: Massard T, Vautrin A, editors. 12th International Conference on Composite Materials (ICCM12), Paris, France, 5–9 July; 1999. [7] Henaff-Gardin C, Rouquie S, Lafarie-Frenot MC, Bertheau D. Fissuration matricielle induite par des contraintes cycliques dans des stratifies croises carbone/epoxyde. Revue des Composites et Materiaux Avances 2001;11–13:383–413. [8] Colin X, Marais C, Cochon JL. Kinetic modelling of weight changes during the isothermal oxidative ageing of bismaleimide matrix. In: DURACOSYS 99, Brussels, Belgium; 1999, p. 49–55. [9] Parvatareddy H, Wang JZ. An evaluation of chemical aging/ oxidation in high performance composites using the Vickers mirco-indentation technique. J Compos Mater 1996;30(2):210–30.
M.C. Lafarie-Frenot, S. Rouquie / Composites Science and Technology 64 (2004) 1725–1735 [10] Bowles KJ. Predicting the durability of PMR-15 composites aged at elevated temperature. In: DURACOSYS 99, Brussels, Belgium, 1999; 441–7. [11] Favre JP, Levadoux H, Ochin T, Cinquin J. Vieillissement des composites a matrice organique aux temperatures moyennes. Un premier bilan. In: Baptiste D, Vautrin A, editors. 10emes Journees Nationales sur les Composites JNC10, Paris, France, AMAC; 1996, p. 205–14. [12] Datoo MH. Residual stresses. In: Mechanics of fibrous composites. Elsevier Applied Science; 1991. p. 367–420. [13] Hancox NL. Thermal effects on polymer matrix composites: Part 1. Thermal cycling. Materials and Design 1998;19(3):85– 91. [14] Henaff-Gardin C, Lafarie-Frenot MC, Brillaud J, ElMahi A. Influence of the stacking sequence on fatigue transverse ply cracking in cross-ply composite laminates. In: Masters JE, editor. Damage detection in composite materials. Philadelphia: ASTM; 1992. p. 236–55. [15] Herakovich CT, Davis JG, Mills JS. Thermal microcracking in Celion 600/PMR-15 graphite/polyimide. In: Thermal stresses in severe environments. New York: Plenum Publishing Corporation; 1980. p. 649–63.
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[16] Rouquie S, Lafarie-Frenot MC. Comparison of damage in carbon/epoxy laminates subjected to thermal fatigue in neutral and oxidative environments. In: ECCM 10, Brugge, Belgium; 2002. [17] Cinquin J, Colombaro AM, et al., Etude et prevision du vieillissement de structures composites a matrice organique soumises a des chargements thermomecaniques repetes. In: Recherche Aeronautique sur le Supersonique, Paris, France, S. Candel, Ministere de la Recherche, Ministere de lÕEquipement, des Transports et du Logement; 2002. p. 125–40. [18] Timmerman JF, Tillman MS, Hayes BS, Seferis JC. Matrix and fibre influences on the cryogenic microcracking of carbon fibre/ epoxy composites. Composites A 2002;33(3):323–9. [19] Pederson CL, Gillespie JW, et al. The effect of isothermal aging on transverse crack development in carbon fiber reinforced cross ply laminates. Polym Compos 1995;16(2):154–60. [20] Lee SH, Nam JD. Thermo-oxidative stability of high performance composites under thermal cycling conditions. J Compos Mater 2001;35(5):433–54. [21] Tsotsis TK, Keller S. Aging of polymeric composite specimens for 5000 hours at elevated pressure and temperature. Compos Sci Technol 2001;61(1):75–86.
Influence of oxidative environments on damage in c/epoxy laminates subjected to thermal cycling M.C. Lafarie-Frenot *, S. Rouquie Laboratoire de Mecanique et de Physique, des Materiaux, – UMR CNRS No. 6617, ENSMA, 1 Av. Clement Ader – B.P. 40109, 86961 Futuroscope-Chasseneuil Cedex, France Received 3 March 2003; received in revised form 12 January 2004; accepted 15 January 2004 Available online 20 February 2004
Abstract The aim of this study is to characterise damage processes in carbon/epoxy laminates submitted to thermal cycling in neutral (nitrogen) and oxidative (air, oxygen) atmospheres. Observations of the polished edges of the specimens by optical microscopy and SEM have revealed the presence of matrix shrinkage and fibre/matrix debondings only on samples tested under oxidative atmospheres. Quantitative analyses of transverse matrix cracking by means of microscopy and X radiography have shown the accelerating effect of an oxidative atmosphere. Moreover, the higher the oxygen concentration, the more significant the acceleration of the cracking development. These results are analysed in terms of matrix oxidation, occurring at the highest temperatures of the cycle, and fatigue matrix cracking, occurring at the lowest ones and due to the prevented differential expansions of the plies. During a thermal cycling test performed in an oxidative atmosphere, a coupling effect between these two damage mechanisms would result in the acceleration of the micro cracking observed. 2004 Elsevier Ltd. All rights reserved. Keywords: B: Durability; B. Matrix cracking; B. Thermomechanical properties; Environmental degradation; B. Fatigue
1. Introduction The present study was conducted in the framework of the French ‘‘Aeronautical Supersonic Research’’ program. The objective of this program is to direct the fundamental research in fields where scientific and technological projections are essential for the design of a new generation of supersonic aircraft. Compared to Concorde, the weight saving required on the aircraft structure is estimated at about 30%. This requirement makes necessary the use of organic composite materials for the manufacturing of most of the structure (Concorde was primarily made of aluminium alloys). This new future supersonic aircraft, designed to fly at Mach 2, will have a lifetime from 3 to 4 times longer than that of the Concorde, corresponding to a minimum of 80,000 h, i.e. approximately 20,000 flights. In this type of *
Corresponding author. Tel.: +33-5-4949-8229; fax: +33-5-49498238. E-mail address: [email protected] (M.C. Lafarie-Frenot). 0266-3538/$ - see front matter 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.compscitech.2004.01.005
application, the materials constitutive of the aircraft structural parts can be subjected to cyclic mechanical loadings and to temperature variations of great amplitude, depending on the phase of the flight, subsonic or supersonic. The objective of a significant reduction in mass, compared to Concorde, would lead to a massive use of composite materials, and particularly of carbon fibre – polymeric matrix laminates. In this context, a better prediction of the long-term durability of composite materials with organic matrix is essential. In particular the study of the coupled effects of mechanical cyclic loads, of temperature variations and of a more or less oxidative environment on the damage of these materials remains to be made. Within the framework of the American program ‘‘High Civil Speed Transport’’, project at Mach 2.4, Arendt et al. [1] described a campaign of tests meant to estimate the durability of some composite materials, candidates to this type of application. Averaged cyclic thermo mechanical conditions are simulated, consisting of the superposition of temperature cycles ()36, 200 C)
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and of stress cycles, the period of which is the mean duration of a flight. This approach leads to very long tests, is very expensive, and requires making a priori the choice of definite materials. On the other hand, Gates [2] proposed another procedure, consisting in the definition of accelerated tests by increasing some critical values of parameters characteristic of physical ageing and chemical degradation of the materials considered. He insists on the need of understanding and characterising the coupling that can appear between these parameters, under the combined effects of a mechanical loading, of the temperature and of the environment. When composite laminates with long continuous fibres are subjected to temperature variations, the mismatch of thermal expansion coefficients of fibres and matrix as well as the difference of ply orientation in the lay up, are such that local stresses appear, which can take part in the degradation of the laminate. When these thermal variations are cyclic, they induce, at the ply level, cyclic stress variations which can be compared, at this scale, to a fatigue phenomenon. Various types of damage similar to those observed in mechanical fatigue result from these cyclic stresses, like transverse matrix cracking, fibre/matrix debonding and delamination [3,4]. When these temperature variations occur in the presence of an oxidative environment, other damaging phenomena appear due to matrix oxidation. Concerning thermoset epoxy polymers, the oxidation phenomenon that occurs during high temperature exposure in oxidative atmospheres, was studied through isothermal ageing experiments by Madhukar et al. [5] and Colin et al. [6] The oxidation of the epoxy matrix results in a loss of mass and a reduction in volume of the matrix, inducing shrinkage of the matrix with respect to the fibres [7]. The study carried out by Colin et al. [8], on epoxy and bismaleimide resins, emphasises for both materials the superficial character of oxidation, which results in the creation of a low thickness oxidised layer on the free edges of the samples subjected to ageing. During ageing, the thickness of this layer grows quickly towards an asymptotic value depending on the studied material and on the temperature of ageing. The detection and the characterisation of this oxidation were also completed by micro hardness Vickers tests [9] on various composite laminates with carbon fibres and different thermoset or thermoplastic matrices. These authors have shown that the oxidised layer has a smaller Vickers hardness compared with that of the virgin material and that this reduction in hardness depends on the atmosphere oxygen concentration. Few works deal with the long-term behaviour of composites subjected to high temperature under oxidative atmosphere. Bowles [10] studied the ageing of T650-35/ PMR15 composites at various temperatures and experimentally found a linear relation between the properties in
compression of the laminates and the thickness of the oxidised layer. The prediction of matrix cracking growth under thermal cyclic variations was not studied as much, as its mechanical fatigue loading counterpart. Favre et al. [11] studied the damage behaviour of various composite materials, with carbon fibres and thermoset matrices (epoxy, cyanate or bismaleimide), subjected to thermal cycling. These authors highlighted an interaction between the matrix micro-cracking of the laminate layers due to the cyclic thermal stresses and oxidation already observed under isothermal ageing. However, as these authors noted, the knowledge of the mechanisms of both crack accumulation and of oxidation is necessary before a relevant prediction can be made. The aim of this experimental study is to characterise the damage growth in carbon/epoxy laminates subjected to cyclic variations of temperature ()50 C/180 C) and to specify the influence of the considered atmosphere: neutral (nitrogen) or oxidative (dry air or pure oxygen). The longer-term goal is to determine protocols of ageing tests, accelerated compared to the tests of thermal cycling presented in this paper. This ambitious objective requires a good preliminary understanding of the oxidation and damage mechanisms, their respective ranges of validity and their interactions.
2. Experimental conditions 2.1. Materials and preparation Thermal cycling tests have been performed on continuous carbon fibre and epoxy matrix composites, selected by EADS in a previous study. The specimens are parallelepipedic, of dimensions 35 · 25 mm2 , the length direction being referred to as the 0 axis (cf. Fig. 1). Four stacking sequences have been investigated: unidirectional [0]8 (UD), cross ply laminates [03 /903 ]S , [453 /)453 ]S , and [45/0/)45/90]S (quasi-isotropic – QI) (their respective thicknesses are 1.04, 1.68, 1.68, and 1.16 mm). The cycle of polymerisation has been optimised by EADS CCR in order to obtain a stable material, to avoid any effect of post cure and change of properties during thermal cycling (cf. Fig. 2).
Fig. 1. Specimen geometry and observations of the polished edges.
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Fig. 2. Polymerisation cycle and post cure of studied composite.
2.2. Methods First, the samples have been smoothly polished on two perpendicular sides. They were examined prior to cycling in order to ensure that there were no initial cracks or defects on the free surfaces. 500 thermal cycles have been prescribed and regularly the samples are removed from the thermal cycling oven to observe damage development: on the free edges of the specimens by optical microscopy and scanning electron microscopy (SEM) and inside by means of penetrant enhanced X radiography. As a precaution the samples were not reused after SEM observations or X radiographs. In the first case, the polished edges were coated with a thin film of gold to avoid any charging effect during SEM observations and, in the second one, zinc iodide solution was used to enhance the contrast of the cracks. In both cases, physical or chemical modifications could induce changes in further damage development. 2.3. Experimental device For this study a specific thermal cycling device has been developed in order to control both the environment and the temperature of the test (Fig. 3).
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Polished and dried samples are placed vertically on racks (cf. Fig. 3(a)), in order to have homogeneous gas flow and temperature distribution around the specimens. Then, they are arranged into a specific enclosure (cf. Fig. 3(b)) having low volume and thermal inertia, allowing constant gas flow with an overpressure of 20 mbar. Various gas mixtures, of controlled relative contents, are obtained from industrial dry gases and pumped into the enclosure. Finally, this device is placed inside a thermal equipment allowing controlled temperature variations to be prescribed (cf. Fig. 3(c)). To assess the influence of oxygen on material ageing, the tests presented here were conducted in three atmospheres: neutral (pure dry nitrogen) and oxidative (dry air reconstituted by 22% O2 + 78% N2 , and pure dry oxygen). The prescribed thermal cycles were triangular, the maximum and minimum temperatures being, respectively, of 180 and )50 C with constant cooling and heating rates of 4 C/min. Fig. 4 shows an example of the real testing thermal cycle prescribed, temperature being measured with a thermocouple embedded in the specimen. It can be noted that the glass transition temperature of this composite material, measured by DSC, has been found around 200–210 C, slightly varying with the material batch considered. In order to accelerate the damage and ageing processes, the maximum temperature of the cycle (180 C) has been chosen higher than the real temperature of the material during a supersonic flight (100 C/130 C) but smaller than Tg. 2.4. Thermo elastic calculations of ply stresses Thermo-elastic calculations [12] give an idea of the level of thermal stresses present in each ply of the laminate for a given temperature. These stresses stem from the mismatch of thermal expansion coefficients of plies with different orientations. The lay up restrains differential thermal deformation of the plies and induces
Fig. 3. Environmentally controlled thermal cycling device.
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Fig. 4. Experimental record of a temperature cycle ()50 C/180 C).
period is equivalent to a sort of thermally driven ‘‘fatigue’’ with DT analogous to the change in stress as previously noted by Hancox [13]. It is important to point out that in an unidirectional laminate, made up of eight plies of same orientation, the thermal stresses are only due to the difference of thermal expansion coefficients between fibre and matrix and not, as for the other laminates of this study, between plies of various orientations.
3. Experimental results 3.1. Damage mechanisms due to thermal cycling
Fig. 5. Thermal cycling and corresponding thermal transverse ply-stresses.
thermal stresses in each oriented layer. Moreover, the magnitude of thermal stresses in the material increases as the temperature of the material deviates from the stress-free temperature. For example, in the case of the thermal cycling between )50 and 180 C, in the different layers of the cross ply laminate [03 /903 ]S , the transverse thermal stresses lie between 0 and 45 MPa. In Fig. 5, the transverse thermal stresses (r22th ) are put in correspondence with the variation of temperature: the cyclic temperature variation induces cyclic transverse stress variations, in each ply of the laminate. This cyclic ‘‘thermal loading’’ with a 2 h
Thermal cycling tests induce various types of damages, depending to a certain extent on the orientation and the thickness of the layers in the lay-up, but mainly on the environment of the specimen. In order to compare damages according to the location of the plies in the stacking sequence of the laminate, microscopic observations have been conducted on two perpendicular sides of the samples. Throughout thermal cycling tests, the observations of the free edges have shown three types of damages: permanent deformation of the matrix due to its shrinkage (Fig. 6(a)), debonding between fibres and matrix (Fig. 6(a)), and matrix cracking (Fig. 6(b)). Considering matrix cracking, two kinds of cracks can be observed. ‘‘Transverse cracks’’ are visible in the plies where the fibre direction is not parallel to the observation plane and, in this study, we call ‘‘transverse cracks’’ the only cracks which cross the whole thickness of the layer concerned (Fig. 6(b)). These cracks are propagating through the matrix but with a direction guided by the fibres; for example, in 90 plies sur-
Fig. 6. Different types of damages. [03 /903 ]S oxygen, 100 cycles.
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rounded by 0 plies, they appear globally perpendicular to the layer interfaces (Fig. 6(b)). Moreover, especially in the thick central layer of the [03 /903 ]S laminates, other kinds of cracks may exist; we called them ‘‘small cracks’’ (Fig. 6(b)). The length of these small cracks can vary between a few fibre diameters to half the thickness of the layer and appear more or less slanting on microscopic observations (Fig. 6(b)). Depending on the experimental conditions (stacking sequence, environment, etc.) and the number of cycles, either none, or some, or all these types of damages could be observed. The damage features as well as their variations, according to the different studied parameters, are described in the following sections. 3.1.1. Matrix shrinkage The example of the [03 /903 ]S cross ply laminate is shown in Fig. 7. In this figure, SEM pictures of the free edges of specimens that have sustained 500 thermal cycles, respectively, in nitrogen, in air and in oxygen, are compared. Note that these pictures have been taken with a 45 tilt of the samples. Comparing the three pictures in Fig. 7, one can see that a significant shrinkage of the matrix only exists on the edges of the specimens tested in oxidative environments. This matrix shrinkage is visualised as dark and deep hollows, imprinted on the initial plane of the polished edge. Moreover, the wider the matrix areas concerned, the deeper are the hollows (Fig. 7). In nitrogen, the very small differences in level observed through the scanning electron microscope between fibres and matrix are similar to those observed on every polished specimen edge before any test. Therefore, the aspect of the edge visualised in Fig. 7(a)) is associated with the polishing stage, whereas the matrix shrinkage shown in Figs. 7(b) and (c) is due to an oxidation process. The matrix shrinkage is observed on the sample surfaces for each laminate tested in oxidative environment (air or oxygen). Whatever the stacking sequence is, these shrinkages appear distributed on each of the free edges. However, deep hollows are systematically observed in rich matrix areas that are present at interfaces between layers of different orientations. The unidirectional case is special because all layers have the same orientation, annihilating the stresses due
Fig. 7. SEM pictures (45 tilt) of the polished edges of [03 /903 ]S laminates subjected to 500 thermal cycles: (a) nitrogen, (b) air and (c) oxygen.
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to the mismatch of thermal expansion between layers of different orientation. However, on such laminates, the matrix shrinkage is also present when the test has been performed in oxidative atmosphere. These residual deformations of the matrix have the same characteristics as the shrinkage observed in the other laminates considered (Fig. 8). 3.1.2. Debonding and crack initiation When matrix shrinkage occurs (i.e. in oxidative environment), the high strain gradients present in matrix areas, close to fibres which have a very high stiffness, lead to high local stresses that can be favourable to debonding between fibres and matrix or to crack initiation as shown in Figs. 9 and 10.
Fig. 8. MEB observation (45 tilt) of the edge of a [0]8 sample after 500 thermal cycles in oxygen.
Fig. 9. QI laminate at 500 cycles in oxygen (45 tilt).
Fig. 10. SEM observations of a [03 /903 ]S specimen subjected to 100 thermal cycles in air.
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As an example, Fig. 9 shows a SEM picture of the edge of the quasi-isotropic laminate [45/0/)45/90]S subjected to 500 thermal cycles in oxygen. In that figure, one can see a deep hole in the matrix, which is located at the interface between the central 90 layer and a 45 ply. In that case, the matrix contraction has induced numerous fibre/matrix debondings in the circumference of the hole. For each test, a first interruption has been made after 100 thermal cycles only. When the tests have been performed in oxidative environment, it appeared that, after 100 cycles, matrix shrinkage is already well developed. Consequently, fibre/matrix debonding, which seems to be connected with this type of damage, has been observed very early during the test. For example, in Fig. 10(a), some small fibre/matrix debondings are present in a [03 /903 ]S specimen subjected to only 100 thermal cycles in air. Moreover, observations of the same specimen, at the same stage of the test, have shown short matrix cracks (Fig. 10(b)). However, one cannot yet be sure that these cracks are initiated from fibre/ matrix debonding, propagating from one such debonding to another? – or if they are induced by another matrix fracture mechanism. 3.1.3. Transverse matrix cracks In Fig. 11, the SEM aspects of transverse cracks present in two [03 /903 ]S specimens after 500 thermal cycles, either in oxygen (Fig. 11(a)) or in nitrogen (Fig. 11(b)), are compared. Once more, important differences in relief of the specimen edges appear: they remain very flat in specimen tested in nitrogen, whereas they are quite uneven in specimen tested in oxygen. These pictures show that the crack opening is larger in O2 than in N2 . In oxygen, the cracks are wide, well opened, allowing observations of the cracked surface of the material: in that case, the fracture pattern appears very brittle, most of the fibres are naked with no traces of residual matrix on them (Fig. 11(a)). Comparatively the cracks propagated in nitrogen are more superficial and narrow (Fig. 11(b)). In that case, the crack opening is so small that the fracture pattern cannot be observed
Fig. 11. SEM observations of transverse ply crack (45 tilt) in [03 /903 ]S subjected to 500 thermal cycles: (a) oxygen and (b) nitrogen.
by SEM and therefore cannot be described from such observations. Equivalent observations of transverse cracks obtained by thermal cycling in air have shown intermediate trends, between those of specimens tested in oxygen and in nitrogen. However, the transverse cracking features (large crack opening, brittle aspect of the fracture, etc.) have been found closer to those of the cracks developed in oxygen: that observation emphasises the influence of the oxidative environment on the transverse cracking mechanisms. To conclude this section concerning microscopic observations of damage due to thermal cycling in neutral or in oxidative environment, the main results can be summarised as following: • Matrix shrinkage is observed on surfaces of specimens tested under oxidative atmosphere only. • In that case, debonding between fibres and matrix are more numerous (and that would be probably related to the above point). • Compared to nitrogen, in oxidative environment (air or oxygen), after a given number of thermal cycles, transverse matrix cracks are more numerous, wider opened, and are increasing in number much faster with numbers of thermal cycles. 3.2. Quantitative analysis of cracking As was shown in the above sections, the damage features as well as the kinetics of its development with number of cycles depend on the specimen environment. In a first step and in order to evaluate more quantitatively these differences, we have focused our study on the sole transverse matrix cracking. Moreover, in this paper, we have chosen to present only the results concerning the cross-ply laminate [03 /903 ]S . Indeed, as well known in fatigue, such a lay-up, with ‘‘thick’’ layers of 3 or 6 plies, is favouring transverse cracking, compared to stacking sequences with thinner layers. The crack measurements (number, length, etc.) are therefore easier and more precise in that laminate. However, in the other stacking sequences considered in this study, the quantitative characterisation of the damage has shown tendencies similar to those emphasised in the [03 /903 ]S laminate and described in the following sections. 3.2.1. X-ray observations The microscopic observations of the polished edges of the samples are complemented by X radiographs which reveal the crack distribution inside the specimen (cf. Figs. 12 and 13). The use of a zinc iodide solution is required to enhance the picture contrast. For crack length measurement, the resolution of this method has been estimated at 1 mm. We have noticed previously that, in the [03 /903 ]S laminate, two types of matrix cracks coexist: ‘‘transverse
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Fig. 12. X radiographs of [03 /903 ]S at 100, 300 and 500 cycles in air.
Fig. 13. Comparison of X radiographs at 500 cycles in nitrogen (a), air (b) and oxygen (c).
cracks’’ which span the entire ply thickness and are wide open, and oblique ‘‘small cracks’’. Although we observed numerous small cracks in the inner layer of that laminate by optical microscopy, none appears on the Xray pictures. This is probably due to their short length, which could be less than the resolution of the X radiographs. In air (cf. Fig. 12), transverse cracking develops significantly between the different observation steps. At 100 cycles, only few short cracks are counted in the central layer. Then, at 300 cycles, these cracks have propagated and increased in number, and transverse cracks have appeared in the external layer, spanning the whole sample length. Finally, at 500 cycles, the transverse crack distribution is quasi-periodic and probably saturated in number. In the central layer some cracks occupy half the width (12.5 mm) but for the most part their length lies between 2 and 10 mm. In the external layers the cracks are more regularly distributed through the sample which appears evenly damaged along its whole length (cf. Fig. 12). In Fig. 13, the X-ray pictures obtained after 500 thermal cycles, in the three studied environments are compared. In nitrogen (Fig. 13(a)), the damage built up at 500 cycles is quite small. Few short cracks (whose length is less than the quarter width of the layer) are seen in the central layer whereas no crack is revealed in the external layers by the microscopic observations. In oxygen the transverse cracking development is similar to that observed in air, but with faster kinetics. In Fig. 13(c) obtained after 500 thermal cycles in oxygen, one can observe that the sample is uniformly damaged, the crack distribution appearing very regular, and more regular than in air (Fig. 13(b)). Moreover, at that step of the thermal cycling, some small grey areas can be seen on X-ray pictures, surrounding some matrix cracks and corresponding to
partial delamination between central and external layers. The data concerning the transverse crack numbers obtained either from the microscopic observation of the polished edges of the samples or from X radiographs have been compared. It appeared that both data are very close when the ‘‘transverse cracks’’ only are considered on microscopic observations (neglecting the ‘‘small cracks’’) and when only the cracks very close to the edge on X radiographs were counted. As a consequence, we have chosen to characterize the damage level by two measurements: the ‘‘transverse crack’’ density measured on the edge of the sample and the total cracked surface measured on the X-ray pictures. 3.2.2. Transverse crack density In Fig. 14, the transverse crack densities measured on the polished edges of the [03 /903 ]S samples cycled in nitrogen, air and oxygen are plotted versus the number of cycles. They are expressed in number of cracks per centimetre, and distinguished according to the position of the considered layer in the stacking sequence: either in the six ply central layer, or in the three ply external layer. In Fig. 14, the full lines correspond to the average development of the crack densities in the central layer, and the dotted lines to that in the external layers. The crack density measurements have been repeated twice, on two different specimens, and for the external layer, each value of crack density plotted in this figure is the average of measurements made on both external layers. It can be seen in Fig. 14 that the scattering on the measurements is rather small, and low enough to consider the average curves as representative of the crack density growth throughout the thermal cycling test. Comparing the three environmental conditions, one can see in Fig. 14 that the transverse cracking development is much faster and more important in an oxidative
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10
oxygen
N cracks/cm
8
air 6
4
2
nit r ogen 0 0
50
100
150
200
250
300
350
400
450
500
Number of cycles Fig. 14. Edge transverse crack densities vs number of thermal cycles ()50, 180 C), [03 /903 ]S , in nitrogen (dark grey), air (thin black line) and oxygen (large black line).
atmosphere (oxygen and air) than in a neutral one (nitrogen). In nitrogen, cracks have been observed in the central layer only, by both optical microscopy and X radiography. There was no crack in the external layers, even after 500 thermal cycles. In that environment, the first cracks appeared between 200 and 300 cycles and they increased in number slightly, up to a level of about 2 cracks/cm at the end of the thermal cycling test. In air, the first cracking onset occurred much earlier than in nitrogen: after 100 thermal cycles, the crack density measured in the central layer is already equal to 3.3 ± 1 cracks/cm. In that layer, the increase in the crack number is fast, up to 300 cycles when saturation is reached with a value of 7.4 ± 0.8 cracks/cm. In the external layers, the cracks initiated later, between 100 and 200 cycles; and their number regularly increased up to a level of 6.9 ± 0.9 cracks/cm at 500 cycles. Considering the data scatter, it appears that the damage level reached at 500 cycles in air is quasi-identical for both central layer and external layers. Nevertheless, two kinds of crack development seem to be involved: a saturation state is reached in the central layer whereas, in the external layers, the increase in the crack number is still fast at the end of the test. We could assume that for a number of cycle higher than 500, the crack density in the external layers is higher than that in the central layer. In oxygen, the kinetics of transverse cracking due to thermal cycling is the fastest we have observed. After only 100 cycles, the crack density in the central layer is 6.1 ± 0.1 cracks/cm, and quasi-saturated, with a saturation value of 7.7 ± 0.6 cracks/cm at 300 cycles. This saturation value can be considered as identical to the one measured in air, in the same central layer. If we examine the cracking development in external layers, we
can assume that the cracks initiated before 100 cycles, since the crack density is equal to 1.6 cracks/cm as early as 100 cycles; then, the number of cracks increased very quickly, up to a crack density of 9.6 ± 0.2 cracks/cm at 500 cycles. The comparison of the transverse crack growth according to the environment of the thermal cycling test emphasises the accelerating effect of an oxidative atmosphere. Moreover, the higher the oxygen concentration, the more significant the acceleration of the damage processes. 3.2.3. Cracked area In order to represent the damage level of each sample, the total cracked areas have been evaluated from X radiograph enlargements. For each crack visualised on an X-ray picture, a cracked area value has been obtained by multiplying its length by the thickness of the layer concerned. The total cracked area is then obtained by summation over all transverse cracks present in each layer of the sample. The cracked area values presented in Table 1 represent the damage levels of [03 /903 ]S specimens (35 · 25 · 1.68 mm3 ), at the end of each test, after 500 thermal cycles. The comparison of the cracked area values according to the test environment points out, very clearly, that the damage level caused by matrix transverse cracking is very low in nitrogen. The trend seen above by measuring Table 1 Cracked area (mm2 ) measured in [03 /903 ]S samples (35 · 25 · 1.68 mm2 ), subjected to 500 thermal cycles ()50 C/180 C)
[03 /903 ]S
Nitrogen
Air
Oxygen
28.5 ± 1
500 ± 18
580 ± 16
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the edge crack densities is enhanced: the cracked surface is higher in oxygen (580 mm2 ) than in air (500 mm2 ) and much larger than in nitrogen (28.5 mm2 ). The relatively small difference of cracked surface values between air and oxygen is due to differences in crack length in the central layer and in crack density in the external layers. Indeed, we have noted in the preceding section that, at 500 cycles, the central layer is saturated with transverse cracks in both oxidative environments: the corresponding cracked area values therefore depend only on the crack lengths, which are smaller in air than in oxygen. Moreover, in the external layers, almost all cracks span the length of the sample, but, as seen in Fig. 14, the crack density at 500 cycles is lower in air than in oxygen. The thermal cycling tests performed on other laminates highlighted the same trends, whatever the stacking sequence. The damage processes are always highly accelerated by an oxidative environment and even more under oxygen than under air. The differences essentially consist in the respective values of crack density and length, depending on the layer thickness, its location in the stacking sequence, etc. For example, we found that the thinner the layer is, the later the first crack appears and for a given cycle number, the shorter the cracks are. These observations are similar to those noted in fatigue of cross-ply laminates with different 90 ply thicknesses [14]. However, in oxidative environments, we found a faster crack development in external layers than that expected by considering their thickness and their location: that could highlight a local influence of the oxidation process. These remarks are similar to those made by Herakovich et al. [15], who studied different graphite/ polyimide lay-ups subjected to various kinds of thermal exposures, and who found that the density of micro cracks depends on fibre orientations, stacking sequence, edge effects and cooling rate. A 2D thermo-elastic calculation of stresses which develop in the layers during the thermal cycling shows that they are identical for the three laminates [03 /903 ]S , [453 /)453 ]S at [45/0/)45/90]S (QI). However, we have noted that, in oxidative environment, the cracking damage induced by 500 thermal cycles are very dependent on the lay-up: in particular, the cracked surface area measured in the [03 /903 ]S laminate (e.g. in oxygen: 580 mm2 ) is twice higher than in the [453 /)453 ]S laminate (210 mm2 ) and much higher than in the QI sample (6.5 mm2 ). The stress calculations take into account only the number of plies of same orientation but neither the thickness of the layers, nor their constraining and nearness to the oxidative environment. Moreover, because of the viscous behaviour of the polymer matrix which is enhanced at high temperature, stress relaxation might be significant over the times involved in these tests. Consequently, the thermo-elastic calculations which only give the level of the initial mean stresses in the plies are inadequate to analyse and discuss the
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experimental results that might depend on the local state of stresses on the free edges of the specimen. Because the oxidation process is related to the diffusion rate of oxygen in the matrix, a local approach might be more efficient to enhance the understanding of the role of the lay-up on the damage development in oxidative atmosphere.
4. Discussion The optical and SEM microscopic observations of the surfaces of samples subjected to thermal cycling ()50 C/180 C), have shown a significant effect of oxygen on the damage mechanisms: it results in a matrix shrinkage, seen on all the samples tested in air and oxygen whereas the surfaces of the specimens tested in nitrogen remain as flat as they were after polishing [16]. The experimental protocol is identical in oxygen, air and nitrogen: same preparation, same time of preliminary drying, same polishing and same thermal cycling. The only parameter changing between the series of tests is the oxygen concentration of the test environment. The differences described above can therefore be associated to this parameter. In oxidative atmospheres, the shrinkage of the matrix and the debonding between fibres and matrix could be induced by the oxidation of the matrix, very active between 120 and 180 C (cf. Fig. 15). Indeed, the superficial oxidation of the epoxy resin induces a reduction in volume [8] resulting in a shrinking of the matrix on the composite surfaces. The significant deformations associated with that shrinkage could induce high stresses on the fibre/matrix interfaces; these stresses can lead to interfacial fracture. Moreover, the oxidised matrix has probably degraded strength properties as well as a lower tenacity. A recent study on oxidised thin films of the same epoxy resin has shown that the failure stress and strain values decrease as the ageing duration increases [17]. As seen on Fig. 10, the zones of matrix shrinkage, characteristic of oxidation phenomena, seem to be favourable sites of debonding between fibres and matrix, and onset of matrix cracking. Temperatur Temperature
oxidation ox 120˚C fa t igue fatigue Time cr acking cracking
Fig. 15. Thermal cycling in oxidative atmosphere and phenomena involved.
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This primer damage mechanism due to oxidation would lead to an easier and earlier crack initiation observed in Fig. 14 when comparing oxygen, air and nitrogen transverse cracking history. The onset and propagation of transverse matrix cracks probably occur at the lowest temperatures of the thermal cycle, where the stresses due to the prevented differential expansions of the plies are most significant. Timmerman et al. [18] have studied various types of carbon/epoxy cross ply laminates [03 /903 ]S performing thermal cycling down to cryogenic temperature. They have observed that, as the temperature decreases, the increase of residual stresses due to the difference of coefficients of thermal expansion between fibres and matrix and between plies of various orientations, are inducing micro cracks in the composite plies. The same remark was made by Pederson and Gillespie [19], who studied isothermal ageing at 177 C of BMI and polyimide unidirectional and cross ply laminates. They show that due to thermal expansion mismatch of the plies, residual ply-stresses present in the cross ply laminate, but absent of the unidirectional specimen, provide the driving force for crack propagation. In neutral environment such as nitrogen, the poor level of cracking observed at the end of the test could only be due to the cyclic variations of ply-stresses induced by the thermal cycling. It could be equivalent to that induced by 500 cycles of ‘‘fatigue’’, with the same ply-stress amplitude. In oxidative environment, this crack growth at the lowest temperature of the thermal cycle could be accelerated by the higher brittleness of the matrix and the interfaces due to oxidation. Moreover, with each onset of a new crack, reactive surface increases, which accelerates the process of local oxidation and damage [20,21]. In that case, a coupling effect between oxidation and matrix cracking would result in an acceleration of the damaging processes.
5. Conclusion This study concerning the damage of carbon/epoxy laminates subjected to cyclic variations of temperature ()50 C/180 C), highlights the influence of the test atmosphere: neutral (nitrogen) or oxidative (air or oxygen) on the damage mechanisms. Observations have shown: • Residual deformations of the matrix, seen as hollows on the sample surfaces and seen only in oxidative atmosphere (air and oxygen). This matrix shrinkage was observed in every lay-up, within the layers as well as at the interface between layers of different orientations. • More numerous and deeper debondings between fibres and matrix in air and oxygen compared to nitrogen.
These two damage mechanisms of the specimen surface could be induced by the matrix oxidation which is very active above 120 C in oxidative atmospheres. These zones, weaken by the oxidation, are favourable sites for crack onset. The observations made during thermal cycling tests have also shown matrix cracking much more developed in oxidative environment than in nitrogen. The coupling between oxidation, due to oxidative atmosphere, and thermo-mechanical cyclic stresses, due to temperature variations, could accelerate the damage process, all the more than the oxygen concentration is high: in that study, the matrix cracking has been found more developed in oxygen than in air.
Acknowledgements The authors acknowledge the financial support of the French Research Department, the French Transport Department, and EADS (CCR Suresnes, France) for supplying the composite plates and samples.
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