Availableonlineatwww.sciencedirect.com DIRECT E噩≈3S SEVIER Journal of the European Ceramic Society 25(2005)599-604 www.elsevier.com/locate/jeurceramsoc Distribution of fibre pullout length and interface shear strength within a single fibre bundle for an orthogonal 3-D woven Si-Ti-C-O fibre/Si-Ti-C-O matrix composite tested at 1100C in air lan J Davies a, * Toshio Ogasawara", Takashi Ishikawa a Dero /dvanced Composite Evaluation Technology Center, Institute of Space Technology and eronautics tralia a department of Mechanical Engineering, Curtin University of Technology. G PO. Box U1987, Perth, WA 6845, Australia Japan Aerospace Exploration Agency (JAXA), 6-13-1 Ohsawa, Mitaka-Shi, Tokyo 181-0015, Japan Received 16 December 2003; received in revised form 10 March 2004; accepted 21 March 2004 Available online 15 June 2004 Abstract The distributions of fibre strength, pullout length, and fibre/matrix interface shear strength within a single fibre bundle were investigated for a 3-D woven SiC/SiC composite tensile tested at 1100C in air. Fibre pullout lengths were largest at the fibre bundle centre with an embrittled region of approximate width 15 um at the perimeter. Whereas the fibre strength varied by less than a factor of 2 across the fibre bundle, the fibre/matrix interface shear strength varied by a factor of -23 with a minimum100= 16 MPa) at the centre and a maximum(. 25+0. 21 GPa) close to the embrittled region. The minimum fibre/matrix interface shear strength required for the transition between pseudo-ductile and brittle behaviour was thus estimated to be 2. 25+0.21 GPa for this composite system C 2004 Elsevier Ltd. All rights reserved Keywords: Composites; Failure analysis; Fibres; Interfaces; SIC/SiC 1. Introduction may increase t to such an extent that crack deflection mech anisms at the fibre/matrix interface are suppressed, + lead The mechanical properties of ceramic matrix composites ing to the formation of an embrittled region characterised by (CMCs)are known to be greatly influenced by the fibre flat fibre fracture surfaces and negligible pullout lengths.3. strength Weibull parameters, So and m, measured in situ A related concern for CMCs containing fibres based on the the composite together with the fibre/matrix interface shear silicon carbide(SiC) system is the formation of an oxide strength, tI Whilst the largest values of tensile strength, o, film at the fibre surface, 5 which acts as a flaw population for CMCs have generally been achieved using low values of and decreases fibre strength. 6 T(typically 1100c)due to sealing of cracks at the based on weakly bonded materials such as pyrolytic carbon specimen surface; 18, 19 oxidation resistance was also found (py-C), boron nitride(BN), 6 and La-monazite(LaPO4).7 to increase with use of thin(<0. 1 um) py-C interfaces 18 However, the extreme oxygen sensitivity of py-C and BN Although several researchers have experimentally in above 500C is a major cause for concern in CMCs that typ- vestigated the effects of oxidation on SiC/SiC composites ically exhibit low-stress matrix microcracking.Oxidation 220-24 and bn2- terraces. in addition of the fibre/matrix interface and exposed fibre surfaces0-13 to other CMCs containing SiC fibres, 8 relatively little data is available for the values of So, m, and t within partially oxidised CMCs, particularly for the variation of these prop- E-mail address: 1.Davies(@curtin.edu. au(IJ. Davies) erties within individual fibre bundles. Such information
Journal of the European Ceramic Society 25 (2005) 599–604 Distribution of fibre pullout length and interface shear strength within a single fibre bundle for an orthogonal 3-D woven Si–Ti–C–O fibre/Si–Ti–C–O matrix composite tested at 1100 ◦C in air Ian J. Davies a,∗, Toshio Ogasawara b, Takashi Ishikawa b a Department of Mechanical Engineering, Curtin University of Technology, G.P.O. Box U1987, Perth, WA 6845, Australia b Advanced Composite Evaluation Technology Center, Institute of Space Technology and Aeronautics (ISTA), Japan Aerospace Exploration Agency (JAXA), 6-13-1 Ohsawa, Mitaka-Shi, Tokyo 181-0015, Japan Received 16 December 2003; received in revised form 10 March 2004; accepted 21 March 2004 Available online 15 June 2004 Abstract The distributions of fibre strength, pullout length, and fibre/matrix interface shear strength within a single fibre bundle were investigated for a 3-D woven SiC/SiC composite tensile tested at 1100 ◦C in air. Fibre pullout lengths were largest at the fibre bundle centre with an embrittled region of approximate width 15m at the perimeter. Whereas the fibre strength varied by less than a factor of 2 across the fibre bundle, the fibre/matrix interface shear strength varied by a factor of ∼23 with a minimum (100±16 MPa) at the centre and a maximum (2.25±0.21 GPa) close to the embrittled region. The minimum fibre/matrix interface shear strength required for the transition between pseudo-ductile and brittle behaviour was thus estimated to be 2.25 ± 0.21 GPa for this composite system. © 2004 Elsevier Ltd. All rights reserved. Keywords: Composites; Failure analysis; Fibres; Interfaces; SiC/SiC 1. Introduction The mechanical properties of ceramic matrix composites (CMCs) are known to be greatly influenced by the fibre strength Weibull parameters, So and m, measured in situ the composite together with the fibre/matrix interface shear strength, τ. 1 Whilst the largest values of tensile strength, σ, for CMCs have generally been achieved using low values of τ (typically <10 MPa2) with subsequently large fibre pullout lengths on the order of several hundred microns,3 recent work has shown τ values of 370 MPa to also be consistent with superior mechanical properties.4 An essential aspect of high strength CMCs is their use of fibre/matrix interfaces based on weakly bonded materials such as pyrolytic carbon (py-C),5 boron nitride (BN),6 and La-monazite (LaPO4).7 However, the extreme oxygen sensitivity of py-C and BN above 500 ◦C is a major cause for concern in CMCs that typically exhibit low-stress matrix microcracking.8 Oxidation of the fibre/matrix interface9 and exposed fibre surfaces10–13 ∗ Corresponding author. E-mail address: I.Davies@curtin.edu.au (I.J. Davies). may increase τ to such an extent that crack deflection mechanisms at the fibre/matrix interface are suppressed,14 leading to the formation of an embrittled region characterised by flat fibre fracture surfaces and negligible pullout lengths.3,15 A related concern for CMCs containing fibres based on the silicon carbide (SiC) system is the formation of an oxide film at the fibre surface,15 which acts as a flaw population and decreases fibre strength.16 Modelling of the oxidation behaviour in CMCs based on the SiC/SiC system has indicated the maximum oxidation rate to occur at intermediate temperatures (e.g. 500–900 ◦C17) with the rate decreasing at higher temperatures (e.g. ≥1100 ◦C18) due to sealing of cracks at the specimen surface;18,19 oxidation resistance was also found to increase with use of thin (≤0.1m) py-C interfaces.18 Although several researchers have experimentally investigated the effects of oxidation on SiC/SiC composites containing py-C2,20–24 and BN25–27 interfaces, in addition to other CMCs containing SiC fibres,28 relatively little data is available for the values of So, m, and τ within partially oxidised CMCs, particularly for the variation of these properties within individual fibre bundles. Such information 0955-2219/$ – see front matter © 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.jeurceramsoc.2004.03.022
1. Davies et al. /Journal of the European Ceramic Sociery 25(2005)599-604 may, for example, shed light on the value of t required for the suppression of crack deflection mechanisms in SiC/SIC composites. Thus, the present work will concern itself with a the investigation of mechanical and physical properties within a single fibre bundle for an orthogonal 3-D woven SiC/SiC composite tensile tested at 1 100oC in air. 2. Experimental procedure The composite under investigation, based on the SiC/SiC system, contained Tyranno LoxM Si-THC-O fibres(800 fibres/bundle)surface-modified so as to achieve a 40 nm arbon-rich layer adjacent to the fibre surface. 29 The aim of 12UA21 5 um surface modification was to promote crack deflection at the carbon-rich layer within the fibre itself, rather than at he fibre surface as is the case for most cmcs. The fibre (b) were woven into an orthogonal 3-d orthogonal configura- tion with fibre volume fractions of.19019 and 0.02 in the x, y, and directions, respectively. Matrix densification of the composite was achieved through the repeated poly mer impregnation and pyrolysis(PIP)of a precursor simi- lar to polytitanocarbosilane(PTCs). The use of similar pre cursors for fibre and matrix components was expected to minimise thermal stresses due to any mismatch in the co- efficients of thermal expansion. This composite system eferred to as"NUSK-CMC'' from the initials of the collab- 3 5 um orating partners and has been the subject of recent work by 321 the authors 29.30 Following machining to a suitable test geometry, 29the Fig. 1. Scanning electron micrographs illustrating the observed fracture composite was heated in air to 1 100C at0.75Cs-, loaded urface types in Tyranno" Si-TiC-O fibres:(a) fracture mirror, and(b) fat and featurele under tension(parallel to the y-axis)to failure, and furnace cooled (initial rate of 3.3Cs); the total time spent at 1 100C being on the order of 600s. Further experimental h, and (iii) whether the fibre exhibited a fracture mirror details are available elsewhere 8,29,31 (Fig. 1(a)) or was flat and featureless(Fig. 1 (b)). A total of In contrast to the complex non-linear stress/strain be- 698 fibres(out of the 800 nominally expected)were mea- aviour noted for similar specimens tested at room temper sured with the difference in number being explained by:(1) ature(RT) and 1200C in vacuum. 8, 29,31 the 1100 C/air shielding of fibres by neighbours, and(ii) the presence of specimens stress/strain curve was approximately linear to holes where fibres had pulled out failure with a low tensile strength(55 MPa)compared to In addition to the above parameters, due to the likely re- the rt(381+41 MPa) and 1200oC/vacuum(405+39 MPa) action of oxygen with the fibre surfaces, So and m were in- cases. In fact, the 1100.C/air specimen tensile strength was vestigated at the centre and edge of the fibre bundle by mea- comparable with that of the stress required for propagation Suring the fracture mirror radius, /'m, of individual fibres and suggesting that the specimen may have failed upon initial fibre strength, 5-2 relationship to determine the individual of matrix cracks within transverse fibre bundles (65 MPa) using the followin matrix microcracking(or shortly thereafter) ollowing failure, the specimen fracture surface was nvestigated using a scanning electron microscope(SEM) (Model JSM-6300F, JEOL, Tokyo, Japan)and the following where Am is known as the "mirror constant and was pre viously determined to be 2.50+0.09MPam /2 for the randomly chosen fibre bundle near the specimen centre: (i SH-Ti-C-O fibres in situ the composite 5.36 and close to position within the fibre bundle, (ii) fibre pullout length the value of 2.51 MPam /2 proposed for nominally similar Nicalon Si-C-O fibres values of S. and m were de duced from cumulative failure probability curves of S after National Aerospace Laboratory of Japan, Ube Industries, Ltd, Shik. applying a suitable correction factor. Fibres wi ibo, Ltd, and Kawasaki Heavy Industries, Ltd are known to be effectively tested at a gauge length, Sc
600 I.J. Davies et al. / Journal of the European Ceramic Society 25 (2005) 599–604 may, for example, shed light on the value of τ required for the suppression of crack deflection mechanisms in SiC/SiC composites. Thus, the present work will concern itself with the investigation of mechanical and physical properties within a single fibre bundle for an orthogonal 3-D woven SiC/SiC composite tensile tested at 1100 ◦C in air. 2. Experimental procedure The composite under investigation, based on the SiC/SiC system, contained Tyranno® LoxM Si–Ti–C–O fibres (800 fibres/bundle) surface-modified so as to achieve a 40 nm carbon-rich layer adjacent to the fibre surface.29 The aim of the surface modification was to promote crack deflection at the carbon-rich layer within the fibre itself, rather than at the fibre surface as is the case for most CMCs. The fibres were woven into an orthogonal 3-D orthogonal configuration with fibre volume fractions of 0.19, 0.19, and 0.02 in the x, y, and z directions, respectively. Matrix densification of the composite was achieved through the repeated polymer impregnation and pyrolysis (PIP) of a precursor similar to polytitanocarbosilane (PTCS). The use of similar precursors for fibre and matrix components was expected to minimise thermal stresses due to any mismatch in the coefficients of thermal expansion. This composite system is referred to as “NUSK-CMC” from the initials of the collaborating partners† and has been the subject of recent work by the authors.29,30 Following machining to a suitable test geometry,29 the composite was heated in air to 1100 ◦C at 0.75 ◦C s−1, loaded under tension (parallel to the y-axis) to failure, and furnace cooled (initial rate of 3.3 ◦C s−1); the total time spent at 1100 ◦C being on the order of 600 s. Further experimental details are available elsewhere.8,29,31 In contrast to the complex non-linear stress/strain behaviour noted for similar specimens tested at room temperature (RT) and 1200 ◦C in vacuum,8,29,31 the 1100 ◦C/air specimen’s stress/strain curve was approximately linear to failure with a low tensile strength (∼55 MPa) compared to the RT (381±41 MPa) and 1200 ◦C/vacuum (405±39 MPa) cases. In fact, the 1100 ◦C/air specimen tensile strength was comparable with that of the stress required for propagation of matrix cracks within transverse fibre bundles (∼65 MPa),8 suggesting that the specimen may have failed upon initial matrix microcracking (or shortly thereafter). Following failure, the specimen fracture surface was investigated using a scanning electron microscope (SEM) (Model JSM-6300F, JEOL, Tokyo, Japan) and the following parameters measured for each fibre visible within a single randomly chosen fibre bundle near the specimen centre: (i) position within the fibre bundle, (ii) fibre pullout length, † National Aerospace Laboratory of Japan, Ube Industries, Ltd., Shikibo, Ltd., and Kawasaki Heavy Industries, Ltd. Fig. 1. Scanning electron micrographs illustrating the observed fracture surface types in Tyranno® Si–Ti–C–O fibres: (a) fracture mirror, and (b) flat and featureless. h, and (iii) whether the fibre exhibited a fracture mirror (Fig. 1(a)) or was flat and featureless (Fig. 1(b)). A total of 698 fibres (out of the 800 nominally expected) were measured with the difference in number being explained by: (i) shielding of fibres by neighbours, and (ii) the presence of holes where fibres had pulled out. In addition to the above parameters, due to the likely reaction of oxygen with the fibre surfaces, So and m were investigated at the centre and edge of the fibre bundle by measuring the fracture mirror radius, rm, of individual fibres and using the following relationship to determine the individual fibre strength, S: 32–34 S = Am √rm (1) where Am is known as the “mirror constant” and was previously determined to be 2.50 ± 0.09 MPa m1/2 for the Si–Ti–C–O fibres in situ the composite35,36 and close to the value of 2.51 MPa m1/2 proposed for nominally similar Nicalon® Si–C–O fibres37 Values of So and m were deduced from cumulative failure probability curves of S after applying a suitable correction factor.38 Fibres within CMCs are known to be effectively tested at a gauge length, δc
1. Davies et al. /Journal of the European Ceramic Society 25(2005)599-604 independent of the composite specimen gauge length and to have the form: 1.38 4(h) rSo λ(m) 1380 oC/vacuum where r is the fibre radius(4.03 um20),(h)is the mean fibre 1200°/ai pullout length, and A(m)is a function only of m. In order 亏04 compare values of So obtained under different conditions, So was normalised to a gauge length, Lo, of 10-m using the Weibull scaling relationship, i.e. 0.0 So(Lo =10-m)=So Fibre pullout length(oqm) Values of t within the fibre bundle were estimated using the Fig 3. Fibre pullout length distributions for an orthogonal 3-D woven following rearrangement of Eq ( 2): 1.38 SiC/SiC composite tested under various conditions 2,20 ri(m)So (4) has been presented in Fig. 2(b) which indicates a strong correlation between fibres with low h values and the ex- istence of flat fracture surfaces(Fig. 1(b); in contrast to 3. Results and discussion fracture mirrors(Fig. 1(a)) which typically occur when crack deflection mechanisms are present. That the major 3.1. Fibre pullout length ity of fibres with significant pullout lengths and fracture mirrors were concentrated towards the centre of the fibre bundle was consistent with previous work2 that suggested maximum oxidation damage to have occurred at the fibre length within the fibre bundle with it being clear that the bundle perimeters. A likely mechanism for this phenomenon majority of fibres, particularly those adjacent to the fibre would be the propagation of matrix cracks in the transverse bundle perimeter, exhibited either zero or negligible fibre fibre bundles, allowing oxygen to access the perimeters of pullout and indicative of brittle failure, i.e. suppression of crack deflection mechanisms at the fibre/matrix interface 14 longitudinal (-axis) fibre bundles due to t being excessively high. Further evidence for this Although the largest h values were observed at the fibre bundle centre. these values were still considerable smaller when compared to those of specimens tested at rT or in vac e in the absence of oxidati age. Fig. 3 com pares fibre pullout length distributions for specimens tested under a variety of conditions.20 with(h)for the 1100.C/air 10, pecimen being an order of magnitude lower compared to that of the RT case(810 um). With reference to Eq(4),it would thus be expected that T for the 1 C/air specimen even in the fibre bundle centre where oxidation damage was minimal, would still be significantly larger compared to the values of 5-10 MPa previously measured in non-oxidised E8E3E 3. 2. Fibre strength parameters Fibre strength distributions(normalised to a gauge length of 10-3m) at the centre and edge2 of individual fibre bun dles in RT and 1 C/air specimens have been presented in Fig. 3 with data for the respective values of So and m being given in Table 1. Whereas the centre and edge fibre strength distributions were similar for the rt case(Fig 4(a)) Fig. 2. Positional dependence of properties within a single fibre bundle in an orthogonal 3-D SiC/SiC composite tested at 1100C in air:(a) 2 For the 1100C/air specimen, fibre strengths fibre pullout length, and(b)existence of flat surface or fracture mirror. could not be measured due to the lack of fracture mirrors(Fig. 2(b). Data Note that fibre pullout lengths in(a) have been shown using a logarithmic for edge fibres was thus calculated using fibres adjacent to the embrittled scale
I.J. Davies et al. / Journal of the European Ceramic Society 25 (2005) 599–604 601 independent of the composite specimen gauge length and to have the form:1,38 δc = 4 h λ(m) = rSo τ (2) where r is the fibre radius (4.03 m20), h is the mean fibre pullout length, and (m) is a function only of m. In order to compare values of So obtained under different conditions, So was normalised to a gauge length, Lo, of 10−3 m using the Weibull scaling relationship, i.e. So(Lo = 10−3m) = So δc 10−3 1/m (3) Values of τ within the fibre bundle were estimated using the following rearrangement of Eq. (2): 1,38 τ = rλ(m)So 4 h (4) 3. Results and discussion 3.1. Fibre pullout length Fig. 2(a) illustrates the spatial dependence of pullout length within the fibre bundle with it being clear that the majority of fibres, particularly those adjacent to the fibre bundle perimeter, exhibited either zero or negligible fibre pullout and indicative of brittle failure, i.e. suppression of crack deflection mechanisms at the fibre/matrix interface14 due to τ being excessively high. Further evidence for this Distance along bundle width ( ∝m) 0 50 100 150 Flat surface Fracture mirror Distance along bundle length (∝m) 0 150 300 450 600 Distance along bundle width ( ∝m) 0 50 100 150 0 ∝m 180 ∝m (a) (b) Fig. 2. Positional dependence of properties within a single fibre bundle in an orthogonal 3-D woven SiC/SiC composite tested at 1100 ◦C in air: (a) fibre pullout length, and (b) existence of flat surface or fracture mirror. Note that fibre pullout lengths in (a) have been shown using a logarithmic scale. 1 10 100 1000 Cumulative failure 0.0 0.2 0.4 0.6 0.8 1.0 Room temperature 1200 oC/vacuum 1300 oC/vacuum 1350 oC/vacuum 1380 oC/vacuum 1100 oC/air 1200 oC/air Fibre pullout length (∝m) Fig. 3. Fibre pullout length distributions for an orthogonal 3-D woven SiC/SiC composite tested under various conditions 2,20. has been presented in Fig. 2(b) which indicates a strong correlation between fibres with low h values and the existence of flat fracture surfaces (Fig. 1(b)); in contrast to fracture mirrors (Fig. 1(a)) which typically occur when crack deflection mechanisms are present. That the majority of fibres with significant pullout lengths and fracture mirrors were concentrated towards the centre of the fibre bundle was consistent with previous work2 that suggested maximum oxidation damage to have occurred at the fibre bundle perimeters. A likely mechanism for this phenomenon would be the propagation of matrix cracks in the transverse fibre bundles,8 allowing oxygen to access the perimeters of longitudinal (y-axis) fibre bundles. Although the largest h values were observed at the fibre bundle centre, these values were still considerable smaller when compared to those of specimens tested at RT or in vacuum, i.e. in the absence of oxidation damage. Fig. 3 compares fibre pullout length distributions for specimens tested under a variety of conditions2,20 with h for the 1100 ◦C/air specimen being an order of magnitude lower compared to that of the RT case (810 m3). With reference to Eq. (4), it would thus be expected that for the 1100 ◦C/air specimen, even in the fibre bundle centre where oxidation damage was minimal, would still be significantly larger compared to the values of 5–10 MPa previously measured in non-oxidised specimens.2 3.2. Fibre strength parameters Fibre strength distributions (normalised to a gauge length of 10−3 m) at the centre and edge2 of individual fibre bundles in RT and 1100 ◦C/air specimens have been presented in Fig. 3 with data for the respective values of So and m being given in Table 1. Whereas the centre and edge fibre strength distributions were similar for the RT case (Fig. 4(a)), 2 For the 1100 ◦C/air specimen, fibre strengths close to the perimeter could not be measured due to the lack of fracture mirrors (Fig. 2(b)). Data for edge fibres was thus calculated using fibres adjacent to the embrittled region.
1. Davies et al./Journal of the European Ceramic Society 25 (2005)599-604 Table I Values of fibre strength Weibull parameters, So and m, measured in situ an orthogonal 3-D woven SiC/SiC composite tested in air Room temperature 100°C/air So(GPa) S0(GPa)(L=10-3m) So(GPa) So (GPa)(Lo 10-m)m 3.04±002 3.78±0. 4.18±0.14 4.13±0.03 225±0.25 2.96±0.09 3.15±0.06 3.93±0. 4.10±0 2.45±0.0 1.39±0.15 8.89±047 a significant difference was noted for the 1 100C/air case populations. As a first approximation, the authors assumed a (Fig. 4(b)) with the normalised So for edge fibres being 62% linear variation in So and m between the centre and furthest that of the centre value. However, even So for the centre fibres with non-zero pullout lengths when applying Eq. (4) fibres(2.25+ 0. 25 GPa)was still considerably below that This assumption was justified on the ground that, as shown of the rt specimen( 3.86 GPa), which itself was approx- later, changes in t from Eq (4) would be dominated by (h) imately 30% below that of the"as received"fibres. The rather than So or n reduction in fibre strength for the 1100 C/air specimen was attributed to the formation of a surface oxide layer on the 3.3. Mean fibre pullout length tributed to increased oxidation and hence a thicker oxide It was next decided to investigate the variation of proper layer. Whereas m for fibres in the RT specimen was approx- ties along the fibre bundle minor axis between perimeter and 8.89+0.47 for edge centre with the data in Fig. 2 bei fibres in the 1100C/air specimen with the increase being at- rical) horizontal strips. Whilst this procedure neglects the tributed to the relatively even thickness of the surface oxide effects of oxidation from either end of the major axis, the layer. The lower m value for centre fibres in the 1100 C/air large major/minor axis length ratio would make oxidation specimen(2.96+0.09)was tentatively attributed to a frac- along the minor axis by far the dominant component Re- tion of the fibres having their strength determined by surface sults presented in Fig. 5 confirm, as previously mentioned flaws(i.e. similar to the rt case)with the remainder being the existence of a strong correlation between fibres exhibit- determined by the surface oxide layer, i.e. two distinct fibre ing pullout and those exhibiting fracture mirror behaviour n fact, the correlation in Fig. 5 is more specific as it clearly Center Edge (a)0.0 0 0. 086 Distance from fibre bundle perimeter(odm) Fig.4. Fibre strength distributions(normalised to a scale length of 10-3m) measured in situ an orthogonal 3-D woven SiC/SiC composite: (a)room Fig. 5. Distribution of fibre properties along the minor axis within an temperature, and (b) 1100C in air. "Centre" and "Edge" refer to the individual fibre bundle in an orthogonal 3-D woven SiC/SiC composite general position within the fibre bundle where the measurements were tested at 1100C in air: (a)mean fibre pullout length, (h), and(b)fraction of fibres exhibiting fracture mirrors
602 I.J. Davies et al. / Journal of the European Ceramic Society 25 (2005) 599–604 Table 1 Values of fibre strength Weibull parameters, So and m, measured in situ an orthogonal 3-D woven SiC/SiC composite tested in air Room temperature 1100 ◦C/air So (GPa) So (GPa) (Lo = 10−3 m) m So (GPa) So (GPa) (Lo = 10−3 m) m Centre 3.04 ± 0.02 3.78 ± 0.13 4.18 ± 0.14 4.13 ± 0.03 2.25 ± 0.25 2.96 ± 0.09 Edge 3.15 ± 0.06 3.93 ± 0.13 4.10 ± 0.05 2.45 ± 0.01 1.39 ± 0.15 8.89 ± 0.47 a significant difference was noted for the 1100 ◦C/air case (Fig. 4(b)) with the normalised So for edge fibres being 62% that of the centre value. However, even So for the centre fibres (2.25 ± 0.25 GPa) was still considerably below that of the RT specimen (∼3.86 GPa), which itself was approximately 30% below that of the “as received” fibres.2 The reduction in fibre strength for the 1100 ◦C/air specimen was attributed to the formation of a surface oxide layer on the fibres;10–13 the lower strength of the edge fibres being attributed to increased oxidation and hence a thicker oxide layer. Whereas m for fibres in the RT specimen was approximately 4.1, the respective value was 8.89 ± 0.47 for edge fibres in the 1100 ◦C/air specimen with the increase being attributed to the relatively even thickness of the surface oxide layer. The lower m value for centre fibres in the 1100 ◦C/air specimen (2.96 ± 0.09) was tentatively attributed to a fraction of the fibres having their strength determined by surface flaws (i.e. similar to the RT case2) with the remainder being determined by the surface oxide layer, i.e. two distinct fibre Fibre strength (GPa) 1 2 3 4 5 6 78 Cumulative failure 0.0 0.2 0.4 0.6 0.8 1.0 Cumulative failure 0.0 0.2 0.4 0.6 0.8 1.0 Center Edge (a) (b) Fig. 4. Fibre strength distributions (normalised to a scale length of 10−3 m) measured in situ an orthogonal 3-D woven SiC/SiC composite: (a) room temperature, and (b) 1100 ◦C in air. “Centre” and “Edge” refer to the general position within the fibre bundle where the measurements were taken. populations. As a first approximation, the authors assumed a linear variation in So and m between the centre and furthest fibres with non-zero pullout lengths when applying Eq. (4). This assumption was justified on the ground that, as shown later, changes in τ from Eq. (4) would be dominated by h rather than So or m. 3.3. Mean fibre pullout length It was next decided to investigate the variation of properties along the fibre bundle minor axis between perimeter and centre with the data in Fig. 2 being divided into (symmetrical) horizontal strips. Whilst this procedure neglects the effects of oxidation from either end of the major axis, the large major/minor axis length ratio would make oxidation along the minor axis by far the dominant component. Results presented in Fig. 5 confirm, as previously mentioned, the existence of a strong correlation between fibres exhibiting pullout and those exhibiting fracture mirror behaviour. In fact, the correlation in Fig. 5 is more specific as it clearly Mean pullout length ( ∝m) 0 20 40 60 80 Distance from fibre bundle perimeter (∝m) 0 10 20 30 40 50 60 Fraction of fibres exhibiting 0.0 0.2 0.4 0.6 0.8 1.0 fracture mirrors (a) (b) Fig. 5. Distribution of fibre properties along the minor axis within an individual fibre bundle in an orthogonal 3-D woven SiC/SiC composite tested at 1100 ◦C in air: (a) mean fibre pullout length, h , and (b) fraction of fibres exhibiting fracture mirrors
1. Davies et al. /Journal of the European Ceramic Society 25(2005)599-604 ↑ 4. Conclusions The distributions of fibre strength, pullout length, and fi- bre/matrix interface shear strength, T, within an individual fibre bundle were investigated for an orthogonal 3-D wo ven SiC/SiC composite tensile tested at 1100C in air. The mean fibre pullout length, (h), varied between 59 um at the centre to zero within an embrittled region(15 um width) 0 at the fibre bundle perimeter. Fibre strength(normalised to a gauge length of 10-m)decreased from 2.25+0.25 GPa Distance from fibre bundle perimeter(om) at the centre to 1.39+0. 15 GPa adjacent to the embrittled Fig. 6. Distribution of fibre/matrix interface shear strength, t, along the region, compared to 3.86+0. 13 GPa for a specimen tensile minor axis within an individual fibre bundle in an orthogonal 3-D wover tested at room temperature. The Sic/SiC composite tested at 1100.C in ai the fibre bundle was 100+ 16 mPa at the centre but this in- creased rapidly to a maximum of 2.25+0. 21 GPa close to indicates the presence of a strong linear relationship between bound for t with respect to suppression of crack deflection the mean fibre pullout length(Fig. 5(a))to the fraction, m, mechanisms at the interface in this composite system.Whilst such a wide variation in t within a single fibre bundle has fm(723±1.2)um not previously been noted, this was attributed to the lack of The value of (h)varied between zero at the perimeter t data available for composites containing short fibre pull- to 59 um at the centre with the lowest non-zero value of out lengths. Overall, the experimental data was consistent (h)being 1.7 um at a distance of 15 um from the bundle with oxygen having surrounded the fibre bundle perimete perimeter; thus providing an estimate for the width of the via matrix cracks in the transverse fibre bundles embrittled region in this specimen, i.e. 15 um. It is clear from Eq (4)that the ratio of (h) between fibre bundle centre and lowest non-zero(h) region would also provide a first Acknowledgements approximation for the ratio of t between these positions (i 9/1.7)and being even higher for the ratio between centre The authors wish to gratefully acknowledge Dr. M and perimeter(where (h)was zero) Shibuya of Ube Industries Ltd, Dr. J. Gotoh of Kawasaki Heavy Industries Ltd, and T. Hirokawa and T. Tanamura of 3.4. Fibre/matrix interface shear strength Shikibo Ltd for the manufacture and supply of all materials used in this study Values of t obtained from Eq (4)have been presented in Fig. 6 with the lowest value of t being 100+ 16 MPa at the fibre bundle centre and approximately 20 times that of References the rT case, illustrating the rapid increase in t which re- sults from even a relatively short exposure time to oxygen I. Curtin, W.A., Theory of mechanical properties of ceramic-matrix at elevated temperature. The value of t gradually increased mposites.J.Am. Ceram. Soc. 1991, 74(11), 2837-2845 2. Davies, 1. J, Ishikawa, T, Shibuya, M. and Hirokawa, T, Fibre away from the bundle centre to reach 149+17 MPa at a dis- strength parameters measured in situ for ceramic-matrix composites tance of 39 um from the bundle perimeter but then increased tested at elevated temperature in vacuum and in air. Compos. Sci. rapidly to a maximum of 2.25+0.21 GPa adjacent to the nbrittled region, with r being necessarily greater than this 3. Davies, L J, Ishikawa, T,Shibuya,M.and Hirokawa,T,Optical within the embrittled region. This value of t at the transition croscopy of a 3-D woven SiC/SiC-based composite. Compos. Sci. echnol.1999,59,429-437 zone between brittle and non-brittle regions thus provides 4. Rebillat, F, Lamon, J, Naslain, R, Lara-Curzio, E, Ferber, M.K. an estimate of the minimum t required for suppression of and Besmann, T. M, Interfacial bond strength in SiC/C/SiC composite crack deflection mechanisms within this composite system materials, as studied by single-fibre push-out tests. J. Am. Ceram. Whilst such high values of t have not previously been re- Soc.1998,81(4),965-978 ported within SiC/SiC composites, the authors attribute this 5. Filipuzzi, L, Camus, G, N: R. and Thebault. J. Oxidation mechanisms and kinetics of lD-SiC/C/SiC composite materials. I.An to the lack of quantitative data available in partially oxidised perimental approach. J. Am. Ceram. Soc. 1994, 77(2), 459-466 CMCs exhibiting h values on the order of several microns 6. Prouhet, S, Camus, G, Labrugere, C Guette, A. and Martin, E, Me However, evidence of high strength fibres exhibiting fracture chanical characterization of Si-C(O) fibre/SiC (CVI) matrix compos mirrors and micron-range pullout lengths, i.e. the require- nterphase. J.Am. Ceram. Soc. 1994, 77(3), 649-656 ments for candidate t values in the gPa range, is available B, Hay, R. S, Marshall, D. B, Morgan, P. E. D. and in the literature, 26, 39 suggesting the data presented in this Influence of interfacial roughness on fibre sliding in oxide omposites with La-monazite interphases J. Am. Ceram. Soc. 2003 work to not be exceptional for CMCs 86(2),305-316
I.J. Davies et al. / Journal of the European Ceramic Society 25 (2005) 599–604 603 Distance from fibre bundle perimeter (∝m) 0 10 20 30 40 50 60 Fibre/matrix interface shear strength (MPa) 0 500 1000 1500 2000 2500 Fig. 6. Distribution of fibre/matrix interface shear strength, τ, along the minor axis within an individual fibre bundle in an orthogonal 3-D woven SiC/SiC composite tested at 1100 ◦C in air. indicates the presence of a strong linear relationship between the mean fibre pullout length (Fig. 5(a)) to the fraction, fm, of fibres exhibiting fracture mirrors (Fig. 5(b)) with h = fm (72.3 ± 1.2)m. The value of h varied between zero at the perimeter to 59 m at the centre with the lowest non-zero value of h being 1.7m at a distance of 15m from the bundle perimeter; thus providing an estimate for the width of the embrittled region in this specimen, i.e. 15 m. It is clear from Eq. (4) that the ratio of h between fibre bundle centre and lowest non-zero h region would also provide a first approximation for the ratio of τ between these positions (i.e. 59/1.7) and being even higher for the ratio between centre and perimeter (where h was zero). 3.4. Fibre/matrix interface shear strength Values of τ obtained from Eq. (4) have been presented in Fig. 6 with the lowest value of τ being 100 ± 16 MPa at the fibre bundle centre and approximately 20 times that of the RT case,2 illustrating the rapid increase in τ which results from even a relatively short exposure time to oxygen at elevated temperature. The value of τ gradually increased away from the bundle centre to reach 149±17 MPa at a distance of 39 m from the bundle perimeter but then increased rapidly to a maximum of 2.25 ± 0.21 GPa adjacent to the embrittled region, with τ being necessarily greater than this within the embrittled region. This value of τ at the transition zone between brittle and non-brittle regions thus provides an estimate of the minimum τ required for suppression of crack deflection mechanisms within this composite system. Whilst such high values of τ have not previously been reported within SiC/SiC composites, the authors attribute this to the lack of quantitative data available in partially oxidised CMCs exhibiting h values on the order of several microns. However, evidence of high strength fibres exhibiting fracture mirrors and micron-range pullout lengths, i.e. the requirements for candidate τ values in the GPa range, is available in the literature,26,39 suggesting the data presented in this work to not be exceptional for CMCs. 4. Conclusions The distributions of fibre strength, pullout length, and fi- bre/matrix interface shear strength, τ, within an individual fibre bundle were investigated for an orthogonal 3-D woven SiC/SiC composite tensile tested at 1100 ◦C in air. The mean fibre pullout length, h , varied between 59 m at the centre to zero within an embrittled region (∼15m width) at the fibre bundle perimeter. Fibre strength (normalised to a gauge length of 10−3 m) decreased from 2.25 ± 0.25 GPa at the centre to 1.39 ± 0.15 GPa adjacent to the embrittled region, compared to 3.86 ± 0.13 GPa for a specimen tensile tested at room temperature. The lowest value of τ within the fibre bundle was 100 ± 16 MPa at the centre but this increased rapidly to a maximum of 2.25 ± 0.21 GPa close to the embrittled region, suggesting this value to be a lower bound for τ with respect to suppression of crack deflection mechanisms at the interface in this composite system. Whilst such a wide variation in τ within a single fibre bundle has not previously been noted, this was attributed to the lack of τ data available for composites containing short fibre pullout lengths. Overall, the experimental data was consistent with oxygen having surrounded the fibre bundle perimeter via matrix cracks in the transverse fibre bundles. Acknowledgements The authors wish to gratefully acknowledge Dr. M. Shibuya of Ube Industries Ltd., Dr. J. Gotoh of Kawasaki Heavy Industries Ltd., and T. Hirokawa and T. Tanamura of Shikibo Ltd. for the manufacture and supply of all materials used in this study. References 1. Curtin, W. A., Theory of mechanical properties of ceramic-matrix composites. J. Am. Ceram. Soc. 1991, 74(11), 2837–2845. 2. Davies, I. J., Ishikawa, T., Shibuya, M. and Hirokawa, T., Fibre strength parameters measured in situ for ceramic-matrix composites tested at elevated temperature in vacuum and in air. Compos. Sci. Technol. 1999, 59, 801–811. 3. Davies, I. J., Ishikawa, T., Shibuya, M. and Hirokawa, T., Optical microscopy of a 3-D woven SiC/SiC-based composite. Compos. Sci. Technol. 1999, 59, 429–437. 4. Rebillat, F., Lamon, J., Naslain, R., Lara-Curzio, E., Ferber, M. K. and Besmann, T. M., Interfacial bond strength in SiC/C/SiC composite materials, as studied by single-fibre push-out tests. J. Am. Ceram. Soc. 1998, 81(4), 965–978. 5. Filipuzzi, L., Camus, G., Naslain, R. and Thebault, J., Oxidation mechanisms and kinetics of 1D-SiC/C/SiC composite materials. I. An experimental approach. J. Am. Ceram. Soc. 1994, 77(2), 459–466. 6. Prouhet, S., Camus, G., Labrugere, C., Guette, A. and Martin, E., Mechanical characterization of Si–C(O) fibre/SiC (CVI) matrix composites with a BN-interphase. J. Am. Ceram. Soc. 1994, 77(3), 649–656. 7. Davis, J. B., Hay, R. S., Marshall, D. B., Morgan, P. E. D. and Sayir, A., Influence of interfacial roughness on fibre sliding in oxide composites with La-monazite interphases. J. Am. Ceram. Soc. 2003, 86(2), 305–316
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