- Email: [email protected]
SiC-based composite creep tested at 1100°C in air
April 2001
Materials Letters 48 Ž2001. 205–209 www.elsevier.comrlocatermatlet
Case study of failure in a glass-sealed SiCrSiC-based composite creep tested at 11008C in air Ian J. Davies a,) , Toshio Ogasawara b, Takashi Ishikawa b, Naoto Suzuki b a
AdÕanced Fibro-Science, Kyoto Institute of Technology, Matsugasaki, Sakyo, Kyoto 606-8585, Japan b Structures DiÕision, National Aerospace Laboratory, Mitaka, Tokyo 181-0015, Japan Received 12 November 1999; received in revised form 4 October 2000; accepted 10 October 2000
Abstract A scanning electron microscopy study was carried out for a glass-sealed SiCrSiC-based composite creep tested at 11008C in air. The specimen fracture surface exhibited two distinct regions; oxidised and unoxidised, while initial oxygen ingression into the specimen was traced to a corner of the specimen. q 2001 Elsevier Science B.V. All rights reserved. Keywords: Ceramic matrix composite; SiCrSiC; Creep; Oxidation protection; Si-Ti-C-O fibre; Glass sealant
1. Introduction It is now well established that ceramic matrix composites ŽCMCs., particularly those containing continuous braided, woven, or knitted fibres, exhibit potential for many applications as high temperature structural materials in the aerospace and space fields w1,2x. However, a continuing problem for CMCs is that of oxidation of the fibrermatrix interface w3x, which tends to suppress crack deflection at the fibrermatrix interface w4,5x and changes the mode of failure from pseudo-ductile to brittle w6,7x. One method to overcome this problem has been to coat w8x or impregnate w9x the CMC with an oxidation barrier material in order to reduce oxygen diffusion into the material and hence provide longer operational lifetimes at elevated temperatures Že.g., )
10008C. in an oxidising atmosphere. One such CMC with an oxidation protection system w10x has been developed in a collaboration between the National Aerospace Laboratory ŽJapan., Ube Industries, Shikibo, and Kawasaki Heavy Industries, and is named ANUSK-CMCB. This composite, based on the SiCrSiC system, has a 3-D woven fibre architecture and has been subject to previous investigation by the authors w10–15x. A recent focus of research has been that of creep behaviour when tested at 11008C to 12008C in air w14,15x. The present work will investigate the failure mechanism for a typical NUSK-CMC specimen that failed during creep testing at 11008C in air and is provided as a case study of failure in a ceramic matrix composite with an oxidation protection system.
2. Experimental procedure )
Corresponding author. Tel.: q81-75-724-7850; fax: q81-75724 7800. E-mail address: [email protected] ŽI.J. Davies..
The composite investigated in this work was based on the SiCrSiC system and contained orthogonal
00167-577Xr01r$ - see front matter q 2001 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 7 - 5 7 7 X Ž 0 0 . 0 0 3 0 4 - 9
206
I.J. DaÕies et al.r Materials Letters 48 (2001) 205–209
3-D woven Tyranno w LoxM Si-Ti-C-O fibres that had been heat-treated in a carbon monoxide atmosphere so as to form a 10 nm SiO x-rich layer at the surface surrounding an inner 40 nm carbon-rich layer w13x. The reasoning behind such a surface treatment was to promote interfacial failure within the carbonrich layer rather than at the actual fibrermatrix interface. Such a procedure thus allows a wider tolerance of bonding strength between fibre and matrix than is normally the case for CMC systems. The total fibre volume fraction was approximately 40% with ; 19% of this being parallel to the loading direction Ži.e., y direction.. Matrix consolidation was achieved through the repeated polymer impregnation and pyrolysis of a precursor similar to polytitanocarbosilane w16x. Prior to mechanical testing, specimens were repeatedly impregnated with a glass sealant based on the SiO 2 –Na 2 O system, which resulted in the filling in of any open porosity together with an additional glass coating at the specimen surface. Details of mechanical testing have been presented elsewhere w14,15x but essentially consisted of uniaxial creep testing at 11008C in air using a specimen gauge length of 30 mm. The actual specimen considered in the present case study investigation failed after a loading time of 2.6 = 10 6 s Ž; 30 days. at 11008C under a creep stress of 140 MPa; equivalent to ; 35% of the monotonic tensile strength at room temperature of 404 MPa. Following failure, the specimen fracture surfaces were carefully removed and investigated using a scanning electron microscope ŽModel JSM-6300F, JEOL, Tokyo, Japan..
3. Results and discussion A general view of the composite fracture surface has been presented in Fig. 1a. The curved black line passing through the specimen was the approximate position of a boundary between two different surface features. The first region was typified by fibre failure characteristics shown in Fig. 2a taken from position ‘A’ in Fig. 1a. This region was characterised by a flat fracture surface with negligible or no fibre pullout. Surfaces of the individual fibres, particularly at the fibre bundle perimeter, were generally smooth ŽFig. 3a. and showed no evidence of Afracture mirrorsB Ži.e., the smooth region adjacent to an initiating
Fig. 1. Fracture surface of a glass-sealed SiCrSiC composite creep tested in air: Ža. scanning electron micrograph, and Žb. schematic representation of oxygen path within the specimen.
defect and surrounded by an area of multiple fracture planes. often observed in ceramic fibres. The lack of substantial fibre pullout, which is a main contributor to the enhanced toughness shown by CMCs, strongly suggested region ‘A’ to have possessed ApoorB mechanical properties. Further evidence for this hypothesis was the presence of flat fibre surfaces that indicated local microcracks to have passed through the fibres without crack deflection at the fibrermatrix interface w4x. This type of fracture feature is characteristic of oxidation damage in CMCs whereupon oxygen ingression has increased the fibrermatrix interface shear strength, t , to such an extent that brittle failure occurs w7x. Indeed, such behaviour had previously been noted by the authors for unsealed
I.J. DaÕies et al.r Materials Letters 48 (2001) 205–209
207
spatial distribution of t had been noted in the composite when tested at room temperature and elevated temperature in vacuum w12x, this phenomenon was attributed to the oxygen path within the composite. Furthermore, the spatial distribution of t suggested that oxidation damage within any particular fibre bundle initiated at the perimeter and moved towards the centre region. This conclusion is in contrast to that of previous researchers w3,17x who suggested the oxygen path at the fibrermatrix interface within fibre bundles to be mainly along the fibre longitudinal axis. One hypothesis consistent with a previous optical microscopy study of unsealed composite w11x is that microcracks existed at the perimeter of the fibre bundle between the fibre bundle and adjoining matrix rich region possibly as a result of the combination of thermal mismatch between the two regions and low fibrermatrix interface shear strength in the pre-tested composite.
Fig. 2. Scanning electron micrographs showing general fracture characteristics observed in a SiCrSiC composite creep tested in air: Ža. AflatB type Žcorresponding to region ‘A’ in Fig. 1a., and Žb. ApulloutB type Žcorresponding to region ‘B’ in Fig. 1a..
composite tensile tested at 11008C in air and attributed to an order of magnitude increase in t w12x. It should be noted, however, that, although difficult to see, fibres at the top half of Fig. 2a Žwhich corresponded to the central longitudinal axis of a fibre bundle. did in fact exhibit fracture mirrors but this was not generally the case for fibre bundles in region ‘A’. The existence of fracture mirrors combined with that of negligible fibre pullout would suggest the fibrermatrix interface shear strength, t , to have been just below the critical value required to suppress crack deflection at the fibrermatrix interface w4x. In contrast to this, the lack of fracture mirrors and fibre pullout for fibres in the lower half of Fig. 2a Žwhich corresponded to the longitudinal perimeter of a fibre bundle. indicated t to have been larger than this critical value in this region. Thus, it can be concluded that t decreased from the perimeter of the fibre bundle towards its centre. As no such
Fig. 3. Scanning electron micrographs showing typical fibre fracture surfaces for SiCrSiC composite creep tested in air: Ža. AflatB type, and Žb. Afracture mirrorB type.
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I.J. DaÕies et al.r Materials Letters 48 (2001) 205–209
In contrast to what has been discussed so far, region ‘B’ in Fig. 1a showed significant fibre pullout within individual fibre bundles ŽFig. 2b., as had also been the case for specimens subject to monotonic tensile loading at room and elevated temperature in vacuum w11x. Such a fracture surface appearance is usually indicative of AgoodB mechanical properties, as fibre pullout behaviour is known to be extremely sensitive to slight increases in t as a result of oxidation damage. However, despite this, it should be noted that some degree of oxidation damage most likely had occurred in the fibre bundle shown in Fig. 2b as the typical fibre pullout length Ž- 100 mm. was significantly smaller than that noted for specimens tested at room temperature and elevated temperature in vacuum Žtypically 300–1000 mm w11x.. The sensitivity of fibre pullout length to t is such that even such a large decrease in fibre pullout length may be attributed to a relatively small increase in t . Fracture surfaces of almost all individual fibres in region ‘B’ showed evidence of fracture mirrors Že.g., at the A2-o’clockB position of the central fibre in Fig. 3b. and this indicated microcrack deflection mechanisms at the fibrermatrix interface to have been operational. Indeed, it was even possible to estimate the strength of the fibre in Fig. 3b to have been ; 2.4 GPa using the size of the fracture mirror radius in Fig. 3b Ž; 1 mm. and fracture mirror constant Ž2.5 = 10 6 MPa m1r2 . w10,12x. From the data so far, it would thus appear that two distinct regions existed within the fracture surface; namely, region ‘A’ that showed significant oxidation damage and region ‘B’ that was relatively undamaged. It was also possible to further conclude that, if region ‘A’ exhibited oxidation damage, then initial ingression of oxygen into the specimen would most likely have taken place somewhere in region ‘A’, such as at the outside surface of the specimen. Closer examination of the specimen outer surface showed the initial ingression point to have been at the top right corner of Fig. 1a, which has been magnified and presented in Fig. 4. The specimen had not been coated prior to SEM examination and hence the non-conducting oxidised regions showed themselves as bright areas compared to the darker unoxidised regions. The arrow in Fig. 4 points to the position where oxygen clearly breached the outer
Fig. 4. Scanning electron micrograph showing the failure origin Žarrow. and approximate glass sealant thickness Žparallel lines. in a SiCrSiC composite creep tested in air.
glass layer Žwhose thickness is indicated by the parallel white lines.. One mechanism that has been suggested to account for initial significant oxygen ingression into the glass-sealed NUSK-CMC specimens is selective evaporation of Na and subsequent crystallisation of the glass sealant which might be expected to result in increased viscosity and a volume contraction of the glass sealant w15x. For a specimen where initial significant oxygen ingression occurred at the corner of the specimen, increased evaporation andror reduced sealant coating thickness at specimen corners might also be expected to play a contributory role w18x. Using information obtained in the present investigation, the predicted oxygen path within the specimen has been represented by arrows in Fig. 1 and it should be noted that oxygen appeared to travel relatively quickly along the fibre bundle longitudinal axes towards the top left corner Ži.e., along the x direction fibre bundles. and bottom right corner Ži.e., along the z direction fibre bundles. in Fig. 1a and b. Such a phenomenon would be consistent with the hypothesis presented earlier that the main route of oxygen into individual fibre bundles was through fibre bundle perimeters where microcracks presumably occurred at the junction between fibre bundle and matrix-rich regions. The existence of such microcracks would inevitably also provide a rapid route for oxygen along the Žperimeter. length of fibre bundles as suggested in Fig. 1.
I.J. DaÕies et al.r Materials Letters 48 (2001) 205–209
4. Conclusions A scanning electron microscopy study of failure was carried out for a glass-sealed SiCrSiC-based specimen that had failed following creep testing for 2.6 = 10 6 s at 11008C in air under a creep stress of 140 MPa. It was concluded that the specimen fracture surface exhibited two distinct regions; oxidised and unoxidised. The oxidised region was characterised by a flat fracture surface with the majority of fibres possessing negligible fibre pullout and no evidence of a fracture mirror. On the other hand, fibres in the unoxidised region showed significant pullout and generally had a fracture mirror. These characteristics were explained in terms of the relative degrees of oxidation damage and increase in the fibrermatrix interface shear strength. The point of initial oxygen ingression into the specimen was traced to one corner of the specimen where oxygen breached the outer glass coating.
References w1x H. Ohnabe, S. Masaki, M. Onozuka, K. Miyahara, T. Sasa, Composites, Part A 30 Ž1999. 489. w2x H. Imuta, J. Gotoh, Key Eng. Mater. vols. 164r165, Trans Tech Publications, Switzerland, 1999, p. 439. w3x R. Naslain, Adv. Compos. Mater. 5 Ž1. Ž1995. 35.
209
w4x H. He, J.W. Hutchinson, Int. J. Solids Struct. 25 Ž9. Ž1989. 1053. w5x A.G. Evans, M.Y. He, J.W. Hutchinson, J. Am. Ceram. Soc. 72 Ž12. Ž1989. 2300. w6x D.A. Woodford, D.R. Van Steele, J.A. Brehm, L.A. Timms, J.E. Palko, JOM Ž1993. 57, May. w7x A.G. Evans, F.W. Zok, R.M. McMeeking, Z.Z. Du, J. Am. Ceram. Soc. 79 Ž9. Ž1996. 2345. w8x H. Hatta, T. Aoki, Y. Kogo, T. Yarii, Composites, Part A 30 Ž1999. 515. w9x J.B. Davis, D.B. Marshall, K.S. Oka, R.M. Housley, P.E.D. Morgan, Composites, Part A 30 Ž1999. 483. w10x I.J. Davies, T. Ishikawa, M. Shibuya, T. Hirokawa, J. Gotoh, Composites, Part A 30 Ž4. Ž1999. 587. w11x I.J. Davies, T. Ishikawa, M. Shibuya, T. Hirokawa, Compos. Sci. Technol. 59 Ž3. Ž1999. 429. w12x I.J. Davies, T. Ishikawa, M. Shibuya, T. Hirokawa, Compos. Sci. Technol. 59 Ž6. Ž1999. 801. w13x T. Ishikawa, K. Bansaku, N. Watanabe, Y. Nomura, M. Shibuya, T. Hirokawa, Compos. Sci. Technol. 58 Ž1998. 51. w14x T. Ishikawa, N. Suzuki, I.J. Davies, M. Shibuya, T. Hirokawa, J. Gotoh, Key Eng. Mater. vols. 164r165, Trans Tech Publications, Switzerland, 1999, p. 197. w15x T. Ogasawara, T. Ishikawa, N. Suzuki, I.J. Davies, M. Suzuki, J. Gotoh, T. Hirokawa, J. Mater. Sci. 35 Ž2000. 785. w16x S. Masaki, K. Moriya, T. Yamamura, M. Shibuya, H. Ohnabe, in: A.G. Evans, R. Naslain ŽEds.., Ceram. Trans. vol. 58, The American Ceramic Society, Westerville, USA, 1995, p. 187. w17x L. Filipuzzi, R. Naslain, J. Am. Ceram. Soc. 77 Ž2. Ž1994. 467. w18x I.J. Davies, T. Ogasawara, T. Ishikawa, N. Suzuki, in: H. Fukuda, T. Ishikawa, Y. Kogo ŽEds.., Proceedings of the ninth US–Japan conference on composite materials, Japan Society for Composite Materials, Tokyo, Japan, 2000, p. 227.
Materials Letters 48 Ž2001. 205–209 www.elsevier.comrlocatermatlet
Case study of failure in a glass-sealed SiCrSiC-based composite creep tested at 11008C in air Ian J. Davies a,) , Toshio Ogasawara b, Takashi Ishikawa b, Naoto Suzuki b a
AdÕanced Fibro-Science, Kyoto Institute of Technology, Matsugasaki, Sakyo, Kyoto 606-8585, Japan b Structures DiÕision, National Aerospace Laboratory, Mitaka, Tokyo 181-0015, Japan Received 12 November 1999; received in revised form 4 October 2000; accepted 10 October 2000
Abstract A scanning electron microscopy study was carried out for a glass-sealed SiCrSiC-based composite creep tested at 11008C in air. The specimen fracture surface exhibited two distinct regions; oxidised and unoxidised, while initial oxygen ingression into the specimen was traced to a corner of the specimen. q 2001 Elsevier Science B.V. All rights reserved. Keywords: Ceramic matrix composite; SiCrSiC; Creep; Oxidation protection; Si-Ti-C-O fibre; Glass sealant
1. Introduction It is now well established that ceramic matrix composites ŽCMCs., particularly those containing continuous braided, woven, or knitted fibres, exhibit potential for many applications as high temperature structural materials in the aerospace and space fields w1,2x. However, a continuing problem for CMCs is that of oxidation of the fibrermatrix interface w3x, which tends to suppress crack deflection at the fibrermatrix interface w4,5x and changes the mode of failure from pseudo-ductile to brittle w6,7x. One method to overcome this problem has been to coat w8x or impregnate w9x the CMC with an oxidation barrier material in order to reduce oxygen diffusion into the material and hence provide longer operational lifetimes at elevated temperatures Že.g., )
10008C. in an oxidising atmosphere. One such CMC with an oxidation protection system w10x has been developed in a collaboration between the National Aerospace Laboratory ŽJapan., Ube Industries, Shikibo, and Kawasaki Heavy Industries, and is named ANUSK-CMCB. This composite, based on the SiCrSiC system, has a 3-D woven fibre architecture and has been subject to previous investigation by the authors w10–15x. A recent focus of research has been that of creep behaviour when tested at 11008C to 12008C in air w14,15x. The present work will investigate the failure mechanism for a typical NUSK-CMC specimen that failed during creep testing at 11008C in air and is provided as a case study of failure in a ceramic matrix composite with an oxidation protection system.
2. Experimental procedure )
Corresponding author. Tel.: q81-75-724-7850; fax: q81-75724 7800. E-mail address: [email protected] ŽI.J. Davies..
The composite investigated in this work was based on the SiCrSiC system and contained orthogonal
00167-577Xr01r$ - see front matter q 2001 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 7 - 5 7 7 X Ž 0 0 . 0 0 3 0 4 - 9
206
I.J. DaÕies et al.r Materials Letters 48 (2001) 205–209
3-D woven Tyranno w LoxM Si-Ti-C-O fibres that had been heat-treated in a carbon monoxide atmosphere so as to form a 10 nm SiO x-rich layer at the surface surrounding an inner 40 nm carbon-rich layer w13x. The reasoning behind such a surface treatment was to promote interfacial failure within the carbonrich layer rather than at the actual fibrermatrix interface. Such a procedure thus allows a wider tolerance of bonding strength between fibre and matrix than is normally the case for CMC systems. The total fibre volume fraction was approximately 40% with ; 19% of this being parallel to the loading direction Ži.e., y direction.. Matrix consolidation was achieved through the repeated polymer impregnation and pyrolysis of a precursor similar to polytitanocarbosilane w16x. Prior to mechanical testing, specimens were repeatedly impregnated with a glass sealant based on the SiO 2 –Na 2 O system, which resulted in the filling in of any open porosity together with an additional glass coating at the specimen surface. Details of mechanical testing have been presented elsewhere w14,15x but essentially consisted of uniaxial creep testing at 11008C in air using a specimen gauge length of 30 mm. The actual specimen considered in the present case study investigation failed after a loading time of 2.6 = 10 6 s Ž; 30 days. at 11008C under a creep stress of 140 MPa; equivalent to ; 35% of the monotonic tensile strength at room temperature of 404 MPa. Following failure, the specimen fracture surfaces were carefully removed and investigated using a scanning electron microscope ŽModel JSM-6300F, JEOL, Tokyo, Japan..
3. Results and discussion A general view of the composite fracture surface has been presented in Fig. 1a. The curved black line passing through the specimen was the approximate position of a boundary between two different surface features. The first region was typified by fibre failure characteristics shown in Fig. 2a taken from position ‘A’ in Fig. 1a. This region was characterised by a flat fracture surface with negligible or no fibre pullout. Surfaces of the individual fibres, particularly at the fibre bundle perimeter, were generally smooth ŽFig. 3a. and showed no evidence of Afracture mirrorsB Ži.e., the smooth region adjacent to an initiating
Fig. 1. Fracture surface of a glass-sealed SiCrSiC composite creep tested in air: Ža. scanning electron micrograph, and Žb. schematic representation of oxygen path within the specimen.
defect and surrounded by an area of multiple fracture planes. often observed in ceramic fibres. The lack of substantial fibre pullout, which is a main contributor to the enhanced toughness shown by CMCs, strongly suggested region ‘A’ to have possessed ApoorB mechanical properties. Further evidence for this hypothesis was the presence of flat fibre surfaces that indicated local microcracks to have passed through the fibres without crack deflection at the fibrermatrix interface w4x. This type of fracture feature is characteristic of oxidation damage in CMCs whereupon oxygen ingression has increased the fibrermatrix interface shear strength, t , to such an extent that brittle failure occurs w7x. Indeed, such behaviour had previously been noted by the authors for unsealed
I.J. DaÕies et al.r Materials Letters 48 (2001) 205–209
207
spatial distribution of t had been noted in the composite when tested at room temperature and elevated temperature in vacuum w12x, this phenomenon was attributed to the oxygen path within the composite. Furthermore, the spatial distribution of t suggested that oxidation damage within any particular fibre bundle initiated at the perimeter and moved towards the centre region. This conclusion is in contrast to that of previous researchers w3,17x who suggested the oxygen path at the fibrermatrix interface within fibre bundles to be mainly along the fibre longitudinal axis. One hypothesis consistent with a previous optical microscopy study of unsealed composite w11x is that microcracks existed at the perimeter of the fibre bundle between the fibre bundle and adjoining matrix rich region possibly as a result of the combination of thermal mismatch between the two regions and low fibrermatrix interface shear strength in the pre-tested composite.
Fig. 2. Scanning electron micrographs showing general fracture characteristics observed in a SiCrSiC composite creep tested in air: Ža. AflatB type Žcorresponding to region ‘A’ in Fig. 1a., and Žb. ApulloutB type Žcorresponding to region ‘B’ in Fig. 1a..
composite tensile tested at 11008C in air and attributed to an order of magnitude increase in t w12x. It should be noted, however, that, although difficult to see, fibres at the top half of Fig. 2a Žwhich corresponded to the central longitudinal axis of a fibre bundle. did in fact exhibit fracture mirrors but this was not generally the case for fibre bundles in region ‘A’. The existence of fracture mirrors combined with that of negligible fibre pullout would suggest the fibrermatrix interface shear strength, t , to have been just below the critical value required to suppress crack deflection at the fibrermatrix interface w4x. In contrast to this, the lack of fracture mirrors and fibre pullout for fibres in the lower half of Fig. 2a Žwhich corresponded to the longitudinal perimeter of a fibre bundle. indicated t to have been larger than this critical value in this region. Thus, it can be concluded that t decreased from the perimeter of the fibre bundle towards its centre. As no such
Fig. 3. Scanning electron micrographs showing typical fibre fracture surfaces for SiCrSiC composite creep tested in air: Ža. AflatB type, and Žb. Afracture mirrorB type.
208
I.J. DaÕies et al.r Materials Letters 48 (2001) 205–209
In contrast to what has been discussed so far, region ‘B’ in Fig. 1a showed significant fibre pullout within individual fibre bundles ŽFig. 2b., as had also been the case for specimens subject to monotonic tensile loading at room and elevated temperature in vacuum w11x. Such a fracture surface appearance is usually indicative of AgoodB mechanical properties, as fibre pullout behaviour is known to be extremely sensitive to slight increases in t as a result of oxidation damage. However, despite this, it should be noted that some degree of oxidation damage most likely had occurred in the fibre bundle shown in Fig. 2b as the typical fibre pullout length Ž- 100 mm. was significantly smaller than that noted for specimens tested at room temperature and elevated temperature in vacuum Žtypically 300–1000 mm w11x.. The sensitivity of fibre pullout length to t is such that even such a large decrease in fibre pullout length may be attributed to a relatively small increase in t . Fracture surfaces of almost all individual fibres in region ‘B’ showed evidence of fracture mirrors Že.g., at the A2-o’clockB position of the central fibre in Fig. 3b. and this indicated microcrack deflection mechanisms at the fibrermatrix interface to have been operational. Indeed, it was even possible to estimate the strength of the fibre in Fig. 3b to have been ; 2.4 GPa using the size of the fracture mirror radius in Fig. 3b Ž; 1 mm. and fracture mirror constant Ž2.5 = 10 6 MPa m1r2 . w10,12x. From the data so far, it would thus appear that two distinct regions existed within the fracture surface; namely, region ‘A’ that showed significant oxidation damage and region ‘B’ that was relatively undamaged. It was also possible to further conclude that, if region ‘A’ exhibited oxidation damage, then initial ingression of oxygen into the specimen would most likely have taken place somewhere in region ‘A’, such as at the outside surface of the specimen. Closer examination of the specimen outer surface showed the initial ingression point to have been at the top right corner of Fig. 1a, which has been magnified and presented in Fig. 4. The specimen had not been coated prior to SEM examination and hence the non-conducting oxidised regions showed themselves as bright areas compared to the darker unoxidised regions. The arrow in Fig. 4 points to the position where oxygen clearly breached the outer
Fig. 4. Scanning electron micrograph showing the failure origin Žarrow. and approximate glass sealant thickness Žparallel lines. in a SiCrSiC composite creep tested in air.
glass layer Žwhose thickness is indicated by the parallel white lines.. One mechanism that has been suggested to account for initial significant oxygen ingression into the glass-sealed NUSK-CMC specimens is selective evaporation of Na and subsequent crystallisation of the glass sealant which might be expected to result in increased viscosity and a volume contraction of the glass sealant w15x. For a specimen where initial significant oxygen ingression occurred at the corner of the specimen, increased evaporation andror reduced sealant coating thickness at specimen corners might also be expected to play a contributory role w18x. Using information obtained in the present investigation, the predicted oxygen path within the specimen has been represented by arrows in Fig. 1 and it should be noted that oxygen appeared to travel relatively quickly along the fibre bundle longitudinal axes towards the top left corner Ži.e., along the x direction fibre bundles. and bottom right corner Ži.e., along the z direction fibre bundles. in Fig. 1a and b. Such a phenomenon would be consistent with the hypothesis presented earlier that the main route of oxygen into individual fibre bundles was through fibre bundle perimeters where microcracks presumably occurred at the junction between fibre bundle and matrix-rich regions. The existence of such microcracks would inevitably also provide a rapid route for oxygen along the Žperimeter. length of fibre bundles as suggested in Fig. 1.
I.J. DaÕies et al.r Materials Letters 48 (2001) 205–209
4. Conclusions A scanning electron microscopy study of failure was carried out for a glass-sealed SiCrSiC-based specimen that had failed following creep testing for 2.6 = 10 6 s at 11008C in air under a creep stress of 140 MPa. It was concluded that the specimen fracture surface exhibited two distinct regions; oxidised and unoxidised. The oxidised region was characterised by a flat fracture surface with the majority of fibres possessing negligible fibre pullout and no evidence of a fracture mirror. On the other hand, fibres in the unoxidised region showed significant pullout and generally had a fracture mirror. These characteristics were explained in terms of the relative degrees of oxidation damage and increase in the fibrermatrix interface shear strength. The point of initial oxygen ingression into the specimen was traced to one corner of the specimen where oxygen breached the outer glass coating.
References w1x H. Ohnabe, S. Masaki, M. Onozuka, K. Miyahara, T. Sasa, Composites, Part A 30 Ž1999. 489. w2x H. Imuta, J. Gotoh, Key Eng. Mater. vols. 164r165, Trans Tech Publications, Switzerland, 1999, p. 439. w3x R. Naslain, Adv. Compos. Mater. 5 Ž1. Ž1995. 35.
209
w4x H. He, J.W. Hutchinson, Int. J. Solids Struct. 25 Ž9. Ž1989. 1053. w5x A.G. Evans, M.Y. He, J.W. Hutchinson, J. Am. Ceram. Soc. 72 Ž12. Ž1989. 2300. w6x D.A. Woodford, D.R. Van Steele, J.A. Brehm, L.A. Timms, J.E. Palko, JOM Ž1993. 57, May. w7x A.G. Evans, F.W. Zok, R.M. McMeeking, Z.Z. Du, J. Am. Ceram. Soc. 79 Ž9. Ž1996. 2345. w8x H. Hatta, T. Aoki, Y. Kogo, T. Yarii, Composites, Part A 30 Ž1999. 515. w9x J.B. Davis, D.B. Marshall, K.S. Oka, R.M. Housley, P.E.D. Morgan, Composites, Part A 30 Ž1999. 483. w10x I.J. Davies, T. Ishikawa, M. Shibuya, T. Hirokawa, J. Gotoh, Composites, Part A 30 Ž4. Ž1999. 587. w11x I.J. Davies, T. Ishikawa, M. Shibuya, T. Hirokawa, Compos. Sci. Technol. 59 Ž3. Ž1999. 429. w12x I.J. Davies, T. Ishikawa, M. Shibuya, T. Hirokawa, Compos. Sci. Technol. 59 Ž6. Ž1999. 801. w13x T. Ishikawa, K. Bansaku, N. Watanabe, Y. Nomura, M. Shibuya, T. Hirokawa, Compos. Sci. Technol. 58 Ž1998. 51. w14x T. Ishikawa, N. Suzuki, I.J. Davies, M. Shibuya, T. Hirokawa, J. Gotoh, Key Eng. Mater. vols. 164r165, Trans Tech Publications, Switzerland, 1999, p. 197. w15x T. Ogasawara, T. Ishikawa, N. Suzuki, I.J. Davies, M. Suzuki, J. Gotoh, T. Hirokawa, J. Mater. Sci. 35 Ž2000. 785. w16x S. Masaki, K. Moriya, T. Yamamura, M. Shibuya, H. Ohnabe, in: A.G. Evans, R. Naslain ŽEds.., Ceram. Trans. vol. 58, The American Ceramic Society, Westerville, USA, 1995, p. 187. w17x L. Filipuzzi, R. Naslain, J. Am. Ceram. Soc. 77 Ž2. Ž1994. 467. w18x I.J. Davies, T. Ogasawara, T. Ishikawa, N. Suzuki, in: H. Fukuda, T. Ishikawa, Y. Kogo ŽEds.., Proceedings of the ninth US–Japan conference on composite materials, Japan Society for Composite Materials, Tokyo, Japan, 2000, p. 227.