Microstructure-property relationships of SiC fibre-reinforced magnesium aluminosilicates—II. Mechanical properties and failure characteristics

Acta mater. Vol. 44, No. I, Pp. 2923-2934. 1996 Copyright 0 1996 Acta Metallurgica Inc. Published by Elsevier Science Ltd

Pergamon 09%7151(95)00387-8

Printed in Great Britain. All rights reserved 1359-6454/96 $15.00 + 0.00

MICROSTRUCTURE-PROPERTY RELATIONSHIPS OF SIC FIBRE-REINFORCED MAGNESIUM ALUMINOSILICATES-II. MECHANICAL PROPERTIES AND FAILURE CHARACTERISTICS A. KUMARt University

of Cambridge,

and K. M. KNOWLES

Department of Materials Science and Metallurgy, Cambridge CB2 342, England

Pembroke

Street,

(Received 31 October 1994; in revised form 6 September 1995) Abstract-Interfacial frictional shear stresses, flexural properties and failure mechanisms are reported for two magnesium aluminosilicates unidirectionally reinforced with Nicalon Sic fibres. Composites A and B were hot-pressed at 1500 and 92O”C, respectively. High values of interfacial frictional shear stresses inferred from Marshall’s analysis of the micro-indentation technique could be attributed in part to the presence of compressive radial stresses at the fibre-matrix interfaces. Although both composites failed non-catastrophically in symmetrical four point bend testing at room temperature, the failure modes were different. Extensive matrix microcracking, fibre failure and then fibre pull-out were commonly observed in composite A. Failure modes in composite B included the formation of a limited number of matrix cracks, the failure of fibres in the matrix crack front and progressive delamination. Our observations demonstrate that the mechanical properties, the interfacial frictional shear stresses and the failure mechanisms of both composites are governed by their microstructural features, in particular the chemistry and structure of the matrix-fibre interfacial region. Copyright 0 1996 Acta Metallurgica Inc.

1. INTRODUCTION It is now well established that the structure and properties of the fibre-matrix interface play a dominant role in controlling the mechanical properties of continuous fibre-reinforced glass and glass-ceramic matrix composites [l]. In Part I of this paper [2], it was demonstrated that the structure and morphology of the interfacial layers in two different magnesium aluminosilicate glass-ceramic matrices unidirectionally reinforced with SIC fibres were quite complex and different. As a result of this, interpretation of the mechanical properties and fracture process is not as straightforward as is often presumed. Several attempts have been made in the past to estimate the mechanical properties of SIC fibrereinforced glass-ceramics. In particular, the Sic/ lithium aluminosilicates (LAS) and Sic/calcium aluminosilicates (CAS) systems have received a lot of attention [3-51. Although the mechanical properties of Sic/magnesium aluminosilicate (MAS) composites have been reported in the literature, attempts have not been made to correlate the properties with the microstructural features of the composites [6-91. The first objective of this paper is to report the mechanical properties of the composites whose

microstructural features have been described in Part I [2]. Estimates of the interfacial frictional shear stress have been made using Marshall’s microindentation technique [lo]. Flexural properties and the failure mechanisms of the composites have been determined using a symmetrical four-point bend test. Finally, the properties and the fracture behaviour of these composites have been correlated with the microstructure of the composites.

TPresent address: Department of Mechanical Engineering, Naval Postgraduate School, Monterey, CA 93943, U.S.A. 2923

2. MATERIALS AND EXPERIMENTAL 2.1. Materials Details of material and processing parameters have been given in Part I of this paper [2]. Briefly, composite A was hot-pressed at 15OO”C, whereas composite B was hot-pressed at 920°C. Composite B was ceramed in air at 1150°C for 1 h, whereas no post-processing heat-treatment was given to composite A. Fibre volume fractions were 0.47 and 0.40 in composites A and B, respectively. 2.2. Experimental 2.2.1. Microhardness and interfacial friction stress. A Vickers diamond on a Leitz microhardness tester was used to indent the fibres for evaluating both the hardness and the interfacial friction stress. A load of 245.25 mN was used to measure the hardness of the fibres in both composites. The total indentation time

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Indenter

I

2b

!

2r

I Matrix’

h

! *

Fibre

Matrix

Fig. 1. Schematic of the indentation method for measurement of matrix-fibre interfacial friction stress (after Ref. [lo]). The relevant parameters have been defined in the text.

was 30 s (15 s for load application and 15 s for dwelling). Diagonals of the indentation impression on the fibres and the matrix were measured on the micrographs recorded using both a scanning electron microscope (SEM, CamscanS2) and a Zeiss optical microscope. The interfacial frictional shear stress was estimated using Marshall’s method [lo]. The fibres were loaded so that indentation impressions were seen in the matrix (Fig. 1). The equations z-

H2a4 n 2ur3E,

(1)

and u =(b

-a)cotti

(2)

derived by Marshall [lo] were used to calculate the interfacial frictional shear stress. In these equations, z is the interfacial frictional shear stress (assumed in this analysis to be constant), 2a is the mean diagonal of the indentation on the fibre, H is the fibre hardness, u is the fibre displacement, 26 is the mean diagonal of the indentation impression on the matrix, 2$(x 148”) is the angle between the opposite edges of the indenter, r is the fibre radius and Ef is the elastic modulus of the fibres. According to the suppliers specifications, the elastic modulus of the Nicalon fibres (NL 202) used to fabricate both composites was x 184 GPa [l 11. This value was therefore used to estimate the interfacial frictional shear stresses via equation (1). However, it should be noted that the values of the elastic modulus of Nicalon fibres reported in the literature vary from 182 to 210 GPa [e.g. 121 and that authors often quote a value of 200 GPa instead [lo, 131. The interface friction stress in equation (1) is inversely proportional to the elastic modulus of the fibre and therefore a relatively small value assumed for the elastic modulus of the fibres will produce a relatively high frictional shear stress and vice-versa. and fractography. 2.2.2. Mechanical testing Samples for flexural testing were cut from composite

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plates using a high speed diamond saw. The as-cut samples were mounted on a well polished aluminium stub specially designed for grinding and polishing large composite samples. All four sides of the samples were first ground with graded SIC grit papers, polished with 6pm diamond paste and finally with 1 pm diamond paste. The composite samples were tested in symmetrical four point bend on a stiff servo-hydraulic machine with a 1 kN load cell to investigate their flexural properties and their fracture behaviour. The tests were carried out with the machine under displacement control. The ramp-rate was 0.025 mm/min. Samples of dimensions approximately 50 x 10 x 0.6 mm and 50 x 5 x 3 mm were used for composites A and B, respectively. The separation between the inner loading points was 16mm for both composites, whereas the separation between the inner and the outer loading points was 10 and 12 mm for composites A and B, respectively. The exact span to depth ratio varied from sample to sample because of small variations in the thicknesses of the samples, but was in the range of 14-16 and 4-5 for composites A and B, respectively. Load-deflection curves were plotted using an x-y recorder. In order to study the damage initiation mechanisms, the test was interrupted at different load levels and the samples were examined ex situ in an optical microscope. A minimum of four samples was tested for both composites. The maximum nominal flexural stress and the elastic modulus were calculated from the linear elastic theory of uniform beams. Fracture surfaces were carefully cut and mounted onto an aluminium stub for examination by SEM. Sample surfaces were gold coated to reduce specimen charging and to enhance contrast in the SEM. The SEM was operated in secondary electron mode.

3. RESULTS

3.1. Microhardness

AND INTERPRETATION

and interface friction

stress

The values of microhardness obtained using a 245.25 mN load were 19-24 and 2632GPa for composites A and B, respectively. An example of micro-indentations used to calculate the microhardness of fibres in both composites is shown for composite B in Fig. 2(a). Debonding at the fibre-matrix interface was observed when a load higher than 245.25 mN was used to indent the fibres in both composites. The fibre hardness obtained experimentally in composite A is lower than that in composite B and also the hardness in both composites is higher than the typical Nicalon fibre hardness (13 GPa) available in the literature [3, lo]. Bleay et al. [13] have recently reported a microhardness value of 8.9 GPa for Nicalon fibres, and although they did not mention the load required to obtain this value, it can be inferred that the load used was less than 0.5 N, the load used in their work to cause the fibres to slide within the matrix.

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It is likely that the low microhardness of fibres in composite A compared with composite B is due to the large-scale diffusion of matrix elements into the fibres in composite A [2], producing a softening effect which reduces the hardness compared with the fibres in composite B. The size of the SIC grains in the SIC-0 Nicalon fibres is on a nanometre level [2], approximately three orders of magnitude smaller than the mean diagonal of the indentations used for estimating the fibre hardness, and so a useful analogy is with the hardness behaviour of glasses as a function of network modifiers [14]. Although the hardness of crystalline ceramics increases generally with a decrease in grain size because dislocations generated by the indenter are blocked by the grain boundaries [15], this effect is clearly not relevant in the present case because of the grain sizes involved. There are three possible reasons for the wide discrepancies reported for the experimentally measured hardness values of Nicalon fibres: (i) an intrinsic variation of hardness between the different batches of fibres; (ii) an indentation-size effect; and (iii) systematic experimental errors. Neither (ii) nor (iii) can account for the differences we have observed between the fibres in composites A and B, because the hardness values were taken at a single load and any systematic errors inherent in the measurement will be the same for both batches. When comparing results from different laboratories, it is important to recognise that in general the microhardness of ceramics is load dependent (known as the indentation-size effect, ISE) and that in general it increases with a decrease in applied load [14].

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Values of ISE indices for a number of crystalline :eramics have been given by Sargent [16] and an :xplanation of the effect in terms of a mixed elasticplastic materials deformation response has been given by Bull et al. [17]. An ISE index, n, for Nicalon fibres IS, to the best of our knowledge, not available, but if data collected by Sargent for hot-pressed SIC, where ‘I z 1.7, can be considered to be a first approximation to what might be expected for the fibres, a simple :alculation combining the equations for Vickers hardness with the ISE force-diagonal power law [14] shows that hardness values H, and H, measured at loads F, and F2, respectively, are related through the equation

n-2

H,

F2 T

HI -=(->

F,

’

and so if F,/F, = 4, H,IH, e 0.784 for n = 1.7. Thus, if different laboratories were to have used loads significantly higher than the ones we have used, some of the discrepancies could be explained at least in part by an ISE. However, this is not in accord with the loads used for measuring fibre hardness reported in the literature [lo, 131. It is therefore more likely that a major reason for the discrepancies in the hardness values of Nicalon fibres lies in the calibration and use of different equipment by various researchers generating systematic measurement errors. To examine this possibility, the hardness of fibres in composite B was measured on indentation equipment at the School of Materials Science, University of Bath, Claverton Down, England. In order to obtain a direct compari-

Fig. 2. Nomarski interference contrast images showing (a, b) indentations obtained on the Leitz microhardness tester in composite B and (b) indentations (indicated by arrows) obtained on the Leco microhardness tester in composite B.

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Fig. 3. Nomarski interference contrast image of a fibre pushed in the matrix in composite son with our own measurements, the hardness of the fibres was measured on the optical microscope on their Leco (M-400) microhardness tester (as in the results quoted by the Bath group [18]) using a Vickers diamond and an indentation load of 245.25 mN. A hardness value of l&13 GPa was obtained. These hardness values are considerably lower than the values obtained using our own Leitz microhardness tester. The indentations made on the two different pieces of equipment are shown in Fig. 2. The indentations made on the Leco arrowed in Fig. 2(b) are representative of those reported elsewhere by the Bath group [19] and are used both for hardness measurements and in the relevant Marshall equation for interfacial frictional shear stress measurements. It should be noted that these indentations are not sharp and that the diagonals are longer than the diagonals obtained on the Leitz using the same load. Moreover, the values of diagonals measured on the optical microscope on the Leco equipment equipped with an higher than the values eye-piece were always measured in a SEM, leading to an underestimate of fibre hardness. Thus, it is clear that, at least for this particular comparison, the considerable differences reported in the apparent fibre hardness on the Leitz and the Leco equipment are due purely to the systematic errors generated when introducing and then measuring indentation diagonals using different equipment. A total load on both the fibre and matrix of 981 mN was sufficient to push-in all the fibres in both composites. A fibre pushed into the matrix in composite A is shown in Fig. 3. Interfacial frictional shear stresses estimated using equation (1) were 2433 and 49-71 MPa in composites A and B, respectively. The values of fibre hardness measured on the Leitz

A.

were used to calculate these values of interfacial frictional shear stresses and the assumption has been made that there is no indentation size effect in Nicalon Sic fibres, as other workers in this area have assumed implicitly. It was observed that most of the fibres at or near the edge of the samples of composite B did not slide in the matrix at 98 1 mN, suggesting higher interfacial frictional shear stresses at or near the edges in comparison with the bulk of the sample. Values of interfacial frictional shear stresses of 136-l 58 MPa were obtained from those fibres which could be pushed-in at or near the edge of the samples. These apparently high shear stresses can arise because of localised oxidation of surface layers during ceraming of composite B in air. Similar trends in results have been reported for a Sic/barium osumilite composite by Bleay and Scott [20]. The interfacial frictional shear stresses in these composites are considerably higher than those reported for Sic/LAS composites [lo, 21-231. This can be attributed at least in part to the differences in residual thermal stresses at the fibre-matrix interfaces. Residual thermal stresses arise from the thermal contraction mismatch between matrix and fibre, and also from the unrelaxed volume changes associated with any phase transformation and crystallisation in the matrix. Brun and Singh [22] have shown that the sliding friction stress is nearly zero when the coefficient of thermal expansion of fibre (a,) is greater than the coefficient of thermal expansion of matrix (a,) and that it increases linearly with the thermal expansion mismatch when tlf< CI,. The coefficient of thermal expansion of LAS is smaller than that of Nicalon fibres and therefore

KUMAR Table 1, Assumed

and materials

Comuosite A B

KNOWLES:

0.47 0.40

parameters Y_

Yr

0.20 0.20

0.15 0.15

Sic

REINFORCED

and estimated

3.1 3.1

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values of residual

thermal

stresses in composites

a, x 10-s (‘C’)

AT (K)

Em (GPaj

Er (GPa)

4.0 4.1

800 800

110 80

184 184

A and B

Qw tMPa)

fl,,m (MPa)

56 45

-28 -27

Y. G(and E refer to Poisson’s ratio. linear thermal expansion coefficient and elastic modulus, respectively. The subscripts, f and m refer to the fibre and matrix respectively. Vr is the volume fraction of the fibres. AT is the difference in temperature between the temperature at which the composite can be assumed to be stress-free (e.g. the glass transition temperature) and room temperature. o,,, is the residual axial stress in the matrix interface in a Lam& coaxial cylinder and Q,,, is the residual radial stress in the matrix at the fibre-matrix model of the composite [24]. The value of a, for composite A is based on the known linear thermal coefficients of Mg0_Al,O,-SiO, glasses and mullite. the relative volume fractions of mullite and glass and reasonable values for the Young’s moduli of the mullite and glass [26], whereas the value of OL,for composite B is given by the suppliers [I l]

residual thermal radial tensile stresses will be present at the fibre-matrix interface in the Nicalon-LAS composites, consistent with low interfacial frictional shear stresses. In contrast, the coefficients of thermal expansion of the matrices in both the composites that we have examined here are higher than that of Nicalon fibres and, therefore, residual radial compressive stresses will be present at the fibre-matrix interface. It should be noted that if either of the composites were to have had only a-cordierite in the matrix, there would have been residual tensile stresses at the fibre-matrix interface, rather than compressive stresses, because the coefficient of thermal expansion of a-cordierite is smaller than that of Nicalon fibres. As explained in Part I [2], the matrices in both the composites consist of cc-cordierite and other phases. Phases with high coefficients of thermal expansion, such as mullite in composite A and enstatite in composite B, help to increase the effective thermal expansion coefficient of the matrix in both cases. An estimation of an upper limit to the residual thermal stresses arising solely from the thermal expansion mismatch can be made using the classical Lame solution of a coaxial fibre and surrounding matrix [24,25]. The relevant equations are given in the Appendix. The residual stresses estimated assuming the various material parameters are given in Table 1. The residual radial stresses in composites A and B at the fibre-matrix interfaces are compressive and comparable in magnitude, -28 and -27 MPa, respectively. If, instead of Marshall’s simple model [lo], the more sophisticated model of the push-down technique of Weihs and Nix [23] is used, in which account is taken not only of a constant sliding resistance term, but also of the residual stress arising from thermal expansion mismatch and a Poisson expansion effect, then the displacement of the top of the fibre, u, for a given load F on the fibre takes the form tl=

(1 - 2vrk) 4

F ~ 2pkzr 1

xln

’ +(z,+~(T,,,)

- r t%l+ P%,,) 2p2k’

II

@F 7cr2

where k

=

Emvi

-w + %I) In these equations, vr and v, are the Poisson’s ratios for the fibre and matrix respectively, E, and Em are the Young’s moduli of the fibre and matrix respectively, Yis the radius of the fibres, p the coefficient of friction at the fibre-matrix interface, orro the residual radial stress at the fibre-matrix interface, and z, is a constant sliding resistance term justified by Weihs and Nix on the basis of remnant fibre surface roughness after fibre-matrix debonding. If we assume that the logarithm term in equation (3) is of the form ln(1 + x) for small x, equation (3) can be rewritten in the form 24%

(1 - 2v,k) Ef

F= 4712r3(7, + pi,,,)’

(5)

The term 2v,k is small in comparison with 1 as it is of the order of vf (which is 0.0225 using the value of 0.15 for vr quoted for Nicalon fibres [21,23]). If we follow Marshall [lo] and let F = 2&H, then to a good approximation, 70 + w-Jr,, =-.

H2a4 n”ur3Ef

This is the same as equation (l), but with 7, + purr0 replacing 7. Coefficients of friction, p, assumed for Sic/LAS composites, in which there is good evidence of a carbon layer at the fibre-matrix interface, are quoted in the range 0.014.32 [21-231. If we therefore assume for SiC/MAS that a value of p = 0.2 is not unreasonable in lieu of any firm experimental data, then a 6 MPa contribution to any “constant” interfacial frictional shear stress would arise from the presence of thermal stresses on these calculations and would be independent of errors arising from, in particular, the measurement of fibre hardness. The values of interfacial frictional shear stresses calculated in equation (1) are very sensitive to the values of fibre hardness and any ISE effect, if present. Thus, if a hardness value of lo-13 GPa were to have been used instead of 2632 GPa for the fibres in composite B in equation (1) the interfacial frictional shear stresses would be reduced to 8-13 MPa before

KUMAR

2928 Table 2. Flexural

Comoosite A R

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properties of composites A and B

Elastic modulus (GPa) 153+ 19 104+4

Nominal ultimate flexural stress WPa) 590* 18 415+69

allowance was made for the effect of any residual compressive stresses at the fibre-matrix interface. This reinforces the doubts that can arise about confidence in the absolute magnitudes of interfacial frictional shear stresses quoted in the literature, and while obviously we have confidence in our own measurements and in the observations of the trends in interfacial frictional shear stress that we have measured, the above analysis shows that we should also quite properly be cautious in the interpretation of interfacial shear stresses in terms of absolute magnitudes. 3.2. Fracture characteristics The flexural properties of both composites are given in Table 2. The average maximum nominal flexural stress (as inferred from the linear elastic theory of uniform beams) and the elastic modulus of composite A are higher than those of composite B. This is, in part, a direct consequence of the higher fibre volume fraction of composite A compared to composite B. Applying the rule of mixtures, the values of the elastic moduli of the matrices lie in the range of 111-130 and 4&50GPa (assuming the fibre modulus is in the range 180-200 GPat) for composites A and B, respectively. These values of the elastic modulus of matrix in composite A are slightly lower than that of mullite (E z 145 GPa, [27]). This is because an appreciable amount of residual glass and small amounts of silica and cordierite were present in the matrix [2], and the elastic moduli of these phases are generally lower than that of mul1ite.J The elastic modulus of the composites depends on elastic moduli of both the fibres and the matrix. Given that composite B was hot-pressed at a much lower temperature than composite A, the elastic modulus of the fibres in the former is not likely to be lower than that in the latter. The elastic modulus of cordierite (z 61-l 17 GPa)§ is lower than the elastic modulus of mullite, and therefore, to a first approximation, the elastic modulus of composite B will be lower than that of composite A. However, the

TAccording to the suppliers specifications, the elastic modulus of Nicalon NL 202 fibres is z 184 GPa [l 11. fThe elastic modulus of Mg0-A&O,-Si02 glasses varies between 103 and 111GPa for a number of compositions [26]. The elastic modulus of vitreous silica is ~73 GPa at room temperature [26]. §The elastic modulus of cordierite glass-ceramics will depend on the type of nucleating agents. The values quoted here are taken from Ref. [27]. Although the reported crystalline phase was orthorhombic cordierite, the composition of the parent glass-ceramic was not mentioned.

ALUMINOSILICATES-II

matrix in composite B also contained enstatite, the elastic modulus of which could not be found in the literature. The elastic modulus of the matrix in composite B calculated by the rule of mixtures is much lower than the elastic modulus of cordierite reported in the literature. It was demonstrated in Part I [2] that composite B contained micro-pores in the matrix due to hot-pressing at a low temperature. In addition to this, some surface delamination cracks were also present in composite B. It is reasonable to assume that these cracks arise from poor interlaminar bonding in these composites. The elastic modulus of monolithic ceramics decreases with an increase in porosity [15,28]. It is therefore most likely that the presence of pores and cracks reduces the elastic modulus of composite B by decreasing the elastic modulus of the matrix (see also Ref. [29]). Typical load-displacement curves for composites A and B are shown in Figs 4(a) and (b), respectively. In both composites, the load increases linearly with displacement at first, and this is then followed by a non-linear regime. The load drops after reaching a peak value. The characteristic difference in the fracture behaviour of the composites is illustrated by the variation in the load after the first load drop. In composite A, the load continues to drop with very small or no build up of load, whereas the load builds up considerably before falling again in composite B. Thus, the damage initiation process (such as matrix cracking or failure in compression) may be the same in both composites, but the mechanisms leading to the final failure are different. Observations during flexural testing of composite A indicated that the first damage always occurred either in tension or in compression, but never in shear. It was difficult to judge whether the damage in compression occurred after the matrix cracking in tension or without any damage on the tension side. If the first damage process involves matrix cracking in the tension side of the flexural beam, subsequent failure in compression will be possible because the actual stresses on the compression side will be higher than those on the tension side (e.g. [3], [30]). Fractographic examination of the samples showed extensive matrix cracking on the tension surface of the flexural beam (Fig. 5). Also seen in Fig. 5 are fibres bridging an opened matrix crack. It can be seen that the fibres have failed away from the plane of the matrix crack. This clearly demonstrates the phenomenon of interface debonding and sliding in the crack wake leading to fibre fracture. The final failure of the composite occurs by propagation of the matrix cracks on the tension side through the thickness and/or by buckling of fibres on the compression side. The fractograph in Fig. 6 shows that the composite sample has failed both in compression and tension. Fibre buckling is seen on the compression side and the fibres pulled-out from the matrix are seen on the tension side of the sample.

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(a)

(b)

250 Stress (MPa)

Stress (MPa)

200’

2””

150 100

100 50

1

I

I

I

I

2

3

Displacement

0.25

0.75

1.0

Displacement

(mm)

nsn

1.25

(mm)

Fig. 4. Typical load-deflection plots of (a) composite A and (b) composite B. Values of nominal stress are indicated based on the linear elastic theory of uniform beams.

Fibre pull-out was commonly observed when the samples were finally separated by hand (Fig. 7). SEM observations revealed smooth fibre surfaces and imprints of fibres on the matrix, suggesting weak bonding between the matrix and the fibre. There are matrix regions that are still attached to fibres (Fig. 7). Therefore, it is possible that the matrix was “blown

Fig. 5. Fractograph

away” during testing as was suggested by matrix dust observed on the flexural testing rig after each test. This may give rise to apparently long fibre pull-out lengths. A systematic study of the fracture process in composite B was carried out by interrupting the test at different loads and examining the samples in an

showing extensive matrix cracking and fibres bridging an opened matrix crack in composite A.

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Fig. 6. Fractograph showing both compression and tension failures in composite A. Fibre buckling is also seen.

optical microscope. It was observed that the first damage always occurred on the tensile side of the beam. However, only a few matrix cracks, perpendicular to the loading plane, were observed on the tension side when the test was stopped at a load corresponding to the onset of non-linear deflection in the load-deflection curve. It is possible that some matrix cracks close upon unloading, but observations after loading beyond the onset of non-linear deflection did not reveal any increase in the crack population. No damage on the compression side of the beam was observed at this stage. Matrix cracks were also seen to have propagated partially through the thickness of the sample. On further loading of the samples, delamination cracks were seen to form and the load dropped abruptly. Fibres were also seen bridging these delamination cracks. This suggests that the delamination cracks in the matrix

Fig. 7. Fractograph

showing

fibre pull-out

crossed over the fibres because of misalignment of the fibres [3 11. The load built up again as the delamination cracks propagated. The load then dropped abruptly, suggesting either the failure of bridging fibres and/or the initiation of further delamination cracks (Fig. 8) in a different laminate. The process of load build-up and drop continued. In some samples, damage on the compression side was also observed at higher loads. In addition to these failure processes, crushing below the loading rollers and shear failure between the inner and the outer loading points were also observed. Shear failure between the outer and inner loading points can be attributed to small span to depth ratios (~4). Only two samples were loaded to the extent that it was possible to separate them by hand. These samples were then used for detailed fractography.

m composite A. Matrix at some places.

debris

attached

to the fibres is seen

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Fig. 8. Fractograph

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showing a matrix crack (indicated by arrows) and the formation of delamination cracks in composite B.

The delamination cracks originated at a matrix crack that had propagated partially through the thickness, as in the example shown arrowed in Fig. 8. Additional matrix cracks were opened perpendicular to the delamination cracks and emanate from the delamination cracks as the delamination cracks propagated [32]. The final failure resulted from the cumulative effect of all these failure processes. The fracture surface of a sample failed by the process of matrix cracking and delamination is shown in Fig. 9(a). The flat region [shown by an arrow in Fig. 9(a)] represents matrix cracking without any obvious fibre pull-out in this region. However, a closer examination of the flat region showed some holes in the matrix, indicating some fibre pull-out; some fibres can also be seen protruding from the surface [Fig. 9(b)]. The pull-out lengths were, however, smaller than 10 pm. The failure of fibres in the matrix crack front resulting in the flat region can be attributed to high interfacial friction stresses at the

edge of the sample compared with the bulk sample (see Section 3.1). Similar fracture behaviour has been reported in the SiC/MAS literature [7,20]. Metcalfe et al. [7] have observed brittle failure and delamination in Sic fibre (Tyranno)-reinforced MAS glass-ceramic composites. They attributed the brittle failure to the non-uniform distribution of fibres, suggesting that the flat regions were relatively devoid of fibres. On the other hand, Bleay and Scott [20] reported that brittle failure in the matrix crack front was due to localised oxidation at the surface of the composites (Nicalon fibre-reinforced barium osumilite) during the heat-treatment. Our experimental results here are in accord with the work of Bleay and Scott [20]. Figure 10 shows a delaminated surface of the sample shown in Fig. 9(a). The bare fibres and the fibre imprints in the matrix clearly indicate poor bonding between the matrix and the fibres. The fibre surfaces appear to be fairly smooth, except for some

Fig. 9. (a) SEM micrograph showing fracture surface of a sample of composite B and (b) fractograph showing fibres protruding from the flat surface shown by an arrow in (a).

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Fig. 10. Fractograph

showing the smooth fibre surfaces and fibre imprints in the matrix. A matrix crack is also seen going around the fibres indicating debonding at weak fibre-matrix interfaces.

small debris attached to the surface. A matrix crack is also seen going around the fibres and this again suggests poor bonding between the fibre and the Figure 11 shows a similar delaminated matrix. surface of the same sample. On this surface a number of broken fibres are seen and the fractured fibre ends are generally flat, indicating tensile brittle failure. Fibres generally fracture at defects such as those shown by arrows in Fig. 11. The fracture process in composite B is similar to the failure of a notched unidirectional SIC fibrereinforced borosilicate glass matrix composite [31]. Bordia et al. [31] observed that the first damage was the formation of a matrix crack (mode-I) at the notch

tip. This mode-1 crack extended through the thickness and was followed by the formation of delamination cracks (mixed-mode I and II) perpendicular to the notch. More mode-1 cracks form on propagation of the delamination crack. Similarly, in composite B, mode-1 matrix cracking starts at the tension surface, and then a matrix crack extends through the thickness, followed by mixed-mode delamination at the matrix crack tip. As the delamination cracks extend, more mode-1 cracks form perpendicular to the delamination cracks. Two important differences in the failure mechanisms of composites A and B are thus: (i) extensive matrix cracking of composite A in

Fig. 11. Fractograph showing the broken fibres in a delaminated surface. The flat fracture surfaces indicate tensile brittle

fracture

at defects

shown

by arrows.

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comparison with composite B; and (ii) fibre pull-out contributing to the work of fracture in composite A, whereas in composite B mixed-mode delamination contributed to toughening. The formation of multiple matrix cracks in brittle matrix composites loaded in tension is a direct consequence of the low failure strain of the matrix compared with that of the reinforcing fibres, provided the fibre volume fraction is more than the critical fibre volume fraction [33]. The formation of only a few matrix cracks in composite B compared to composite A is due mainly to the failure of fibres in the matrix crack front at the tension side of the samples of composite B. The localised oxidation of fibres at the surface of composite plates during ceraming makes the composite notch-sensitive, i.e. the propagation of a matrix crack through the thickness is accompanied by failure of fibres at the crack tip instead of debonding at the fibre-matrix interface. Delamination in composite B was because of poor interlaminar bonding arising from the relatively low consolidation temperature. The lower matrix-fibre interfacial shear stresses in the interior of the composite samples, protected from the atmosphere during ceraming, will also facilitate crack-defection. This was confirmed by an examination of the as-received samples which showed delamination cracks. It is therefore obvious that the inferior mechanical properties of composite B are mainly due to the hot-pressing at low temperatures which results in rather porous and incompletely consolidated composites.

4.

DISCUSSION

AND CONCLUSIONS

In this paper, an extensive investigation of the fracture behaviour of Nicalon fibre-reinforced magnesium aluminosilicate matrix composites has been conducted. The as-measured values of fibre hardness in both composites are higher than those normally reported in the literature. The fibre hardnesses are reproducibly lower in the composite consolidated at the higher temperature. It has been demonstrated by a comparison of test procedure in our laboratory and another that at least some of the large variation in the reported values of fibre hardness in the literature can be attributed to the calibration and use of different equipment used by various researchers generating systematic measurement errors. The maximum flexural stress and the elastic modulus of composite A were higher than those of composite B. This was because of the higher fibre volume fraction of composite A. It is also suggested that the presence of pores and interlaminar cracks in composite B reduces the elastic modulus of the matrix, thereby further reducing the elastic modulus of composite B. The strength and modulus values for both composites are higher than those reported previously for similar composites [6, 34, 351.

ALUMINOSILICATES--II

2933

Composite A failed either in tension or in compression but never in shear. Extensive matrix cracking was commonly observed in composite A. Composite B failed by a combination of mode-1 matrix cracking and mixed-mode delamination processes. Although both the composites failed noncatastrophically, the toughening mechanisms were different. Fibre pull-out was the principal toughening mechanism for composite A, whereas progressive delamination was the main toughening mechanism for composite B. It has been demonstrated by a consideration of the coefficients of thermal expansions of the various phases present in the two composites, that residual compressive thermal stresses were present in both composites at the fibre-matrix interfaces. Such stresses can increase the apparent interfacial frictional shear stresses significantly, as a consideration of the model proposed by Weihs and Nix [23] was able to show. Despite the presence of apparently high interfacial frictional shear stresses in composite A, debonding and fibre pull-out occurred during flexural testing. Hence, high interfacial frictional shear stresses inferred from Marshall’s [lo] micro-indentation technique do not necessarily imply a strong bond between the fibre and the matrix. This is in agreement with the presence of the turbostratic carbon layer at the fibre-matrix interface [2], which is able to facilitate debonding at the fibre-matrix interface and eventually cause fibre pull-out at sufficiently high loads. However, the presence of a carbon-rich interlayer does not always lead to fibre pull-out, as is the case for composite B. Other factors, such as surface oxidation during ceraming heat-treatments, can induce “notch-sensitivity”. A complex fracture process involving propagation of matrix cracks through the thickness and then progressive delamination can still impart some damage tolerance. The full potential of composite B could not be exploited due to poor interlaminar bonding and also due to surface oxidation during ceraming heat-treatment. Acknowledgements-We would like to thank Pilkington plc for the provision of samples of fibre-reinforced glassceramics. AK would like to thank the Cambridge Commonwealth Trust for financial support. REFERENCES

1. M. H. Lewis and V. S. R. Murthy. Comp. Sci. Technol. 42, 221 (1991). 2. A. Kumar and K. M. Knowles, Acta. mater. 44, 2901

(1996). 3. D. B. Marshall and A. G. Evans, J. Am. Ceram. Sot. 68, 225 (1985). 4. K. M. Prewo, J. Murev. Sci. 21, 3590 (1986). 5. D. S. Beyerle, S. M. Spearing, F. W. Zok and A. G. Evans, J. Am. Ceram. Sot. 75, 2719 (1992). 6. M. Y. Chen. J. M. Battison and Tai-I1 Mah, J. Mater. Sci. 24, 3213 (1989). 7. B. L. Metcalfe, I. W. Donald and D. J. Bradley, J. Mater. Sci. 27, 3075 (1992).

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8. A. M. Daniel and M. H. Lewis, Ceram. Engng Sci. Proc. 14, 131 (1993). 9. A. Chamberlain, M. W. Pharaoh and M. H. Lewis, Ceram. Emma 14. 939 (1993). 10. D. B. Marshall, i. Am.‘Ceram. Sot. 67, C-259 (1984). 11. J. Ridealgh, Pilkington Technology Centre, Lancashire, UK, private communication. 12. J. Persh, Ceram. Engng Sci. Proc. 9, 529 (1988). 13. S. M. Bleay, V. D. Scott, B. Harris, R. G. Cooke and F. A. Habib, J. Mater. Sci. 27, 2811 (1992). 14. I. J. McColm, Ceramic Hardness. Plenum Press, New York (1990). 15. R. W. Rice, Treatise Mater. Sci. Technol. 11,200 (1977). 16. P. M. Sargent, in Microindentation Techniques in Materials Science and Engineering, ASTM STP 889 (edited by P. J. Blau and B. R. Lawn), p. 160. American Society for Testing and Materials, Philadelphia PA (1986). 17. S. J. Bull, T. F. Page and E. H. Yoffe, Phil. Mag. Lett. 59, 281 (1989). 18. S. M. Bleay, private communication. 19. S. M. Bleay and V. D. Scott, J. Mater. Sci. 26, 2229 (1991). 20. S. M. Bleay and V. D. Scott, J. Mater. Sci. 27, 825 (1992). 21. D. K. Shetty, J. Am. Ceram. Sot. 71, C-107 (1988). 22. M. K. Brun and R. N. Singh, Adv. Ceram. Mater. 3,506 (1988). 23. T. P. Weihs and W. D. Nix, J. Am. Ceram. Sot. 74,524 (1991). 24. L. N. McCartney, Proc. R. Sot. Lond. A425,215 (1989). 25. C. M. Warwick and T. W. Clyne, J. Mater. Sci. 26,3817 (1991). 26. N. P. Bansal and R. H. Doremus, Handbook of Glass Properties. Academic Press, London (1986). 27. Ceramic Source, Vol. 6, Published by the American Ceramic Sot. (1990-91). 28. R. W. Davidge, Mechanical Behaviour of Ceramics. Cambridge University Press, Cambridge (1979). 29. D. C. Phillios. R. A. J. Sambell and D. H. Bowen. J. Mater. SC>. 7, 1454 (1972). 30. S. Jansson and F. A. Leckie, Acta metall. mater. 40, 2967 (1992). 31. R. K. Bordia, B. J. Dalgleish, P. G. Charalambides and A. G. Evans, J. Am. Ceram. Sot. 74, 2176 (1991). 32. A. Kumar, Ph.D. thesis, University of Cambridge (1994). 33. J. Aveston, G. A. Cooper and A. Kelly, Proc. Conf on

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Properties of Fibre Composites, p. 15 IPC Science and Technology Press, NPL, Guildford, UK (1971). 34. B. A. Bender, D. Lewis, W. S. Coblenz and R. W. Rice, Ceram. Engng Sci. Proc. 5, 513 (1984). 35. R. Chaim, D. G. Brandon and L. Baum, Ceram. Engng Sci. Proc. 9, 695 (1988).

APPENDIX The model used here is the classical Lame “shrink-fit” model of a cylinder of fibre surrounded by, and bonded to, a cylinder of matrix. The system is presumed to be stress-free initially in the absence of any externally applied stresses. The composite cylinder is then subjected to a uniform temperature change T relative to this initial state, with the fibre and matrix remaining fully bonded. The residual thermal stresses in the matrix arising from the differences between the coefficients of thermal expansion of fibre and matrix in a composite are calculated from the following equations [24]

u,+y,1)

(A.2)

e=a,T 1 a,=E{VfEfc++ E + 2lV,

(A.3)

V,,,E,,,am Vr(v, - v&K1 + %,)a, - (1 +

vfbfl)

E, = V,E, + V, E,,, + 212V,,, I’&,, - v$ 1

1 + v,

n=E,+

(1 - 2v,)(l ElII

(A.4) (A.5)

+ v,)Vr + (1 - 2vr)(l + vr)Vnl

(A,6)

Er

dJ 1 _=(v,-vr)s+(l+vr)a,T-(l+v,)a,T.

(A.7) R2 0 1 In these equations the subscripts m, f and c refer to matrix, fibre and composite, respectively. c, and o, are the residual axial (in the matrix) and radial (at the fibre-matrix interface) thermal stresses. E, V, R, T, a, E and v refer to elastic modulus, volume fraction, fibre radius, temperature change, coefficient of thermal expansion, thermal strain and Poisson’s ratio, respectively.