COMPOSITES SCIENCE AND TECHNOLOGY ELSEVIER Composites Science and Technology 59(1999)801-811 Fibre strength parameters measured in situ for ceramic-matrix composites tested at elevated temperature in vacuum and in air lan J Davies Takashi Ishikawa, * Masaki Shibuya, Tetsuro Hirokawa rframe Division, National Aerospace Laboratory, 6-13-1 Ohsawa, Mitaka-Shi, Tokyo 181, Japan cOrporate Research and Development, Ube Industries Ltd, 1978-10 Kogushi, UbeShi, Yamaguchi 755, Japan Research and Development Department, Industrial Textile Division, Shikibo Ltd, 1500-5 Shibahara-Minami, Yokaichi-Shi, Shiga 527, Japan Received 3 July 1997; received in revised form 15 July 1998; accepted 7 January 1999 Abstract In situ fibre fracture characteristics have been investigated for Si-Ti-C-O fibres after tensile testing up to 1380"C in vacuum and in air. Specimens tested in air at 1 100 and 1200 C generally had flat fracture surfaces with less than 20% of fibres exhibiting fracture mirrors: this is attributed to oxygen ingress into the fibre bundles. Fibre strength characteristics normalised to a 10-3 m gauge length indicated that fibres tested in air at elevated temperature have significantly lower strengths and average Weibull parameter. m, compared to the room-temperature, 1200 and 1300 C/vacuum cases, and this is attributed to oxygen damage of the fibre toge- her with oxidation of the fibre/matrix interface. The fibre/matrix interface shear strength, t, was low for the room-temperature specimens and increased slightly with temperature when tested in vacuum, possibly as a result of a change in the thermal mismatch between fibres and matrix. Values of r for specimens tested at 1100 and 1200 C in air were an order of magnitude greater than those for room-temperature specimens, indicating a significant degree of oxidation damage at the fibre/ matrix interface to have occurred C 1999 Elsevier Science Ltd. All rights reserved 1. ntroduction shear strength, t, and fibre properties such as the radius r, and the Weibull strength parameters, So and m omposites(CMCs) that utilise con- Values for So and m may be obtained by measuring the tinuous fibres as reinforcement have many potential in situ strength of fibres and fitting a two-parameter high-temperature structural applications, particularly in Weibull curve [7] to the resulting data,i.e the area of space re-entry vehicles. Recent advances include production of 3-D woven composites based on F=1-c-(÷) the sic/Sic system that possess short-term tensile (1) strengths of nearly 400 MPa in vacuum at room tem- perature and 1200C with tensile strains to failure in where F is the cumulative failure probability of fibres at excess of 1%[1-3]. However, the mechanical properties a stress, S, and S, and m are empirical constants known of CMCs often degrade at elevated temperature in the as the Weibull strength parameters. Although the ex situ presence of oxygen which is known to attack the fibre/ strength of ceramic fibres is often well known, there is a matrix interface. Although studies have shown that relative dearth of data concerning in situ fibre strength sealing the surface of CMCs with a glass-based com- One method of estimating in situ fibre strength involves pound may allow similar mechanical properties to be the measurement of mirror radi-a feature often pre- achieved at elevated temperature in air and in vacuum sent on the fracture surface of ceramic fibres. Fig. I [4], improvement of the fibre/matrix interface is still a illustrates such a fracture mirror observed on the frac- ture surface of a Tyranno Si-Ti-C-O fibre It can be Recent work [5,6] has indicated that several impor seen that the fracture mirror is cent tant composite mechanical properties may be predicted to the initiating defect in the fibre and is surrounded by from a knowledge only of the fibre/matrix interface a region of multiple fracture planes. Past research has shown most fracture mirrors to 4 Corresponding author. initiate at flaws present on the surface of the fibre [8, 9 0266-3538/99/S- see front matter C 1999 Elsevier Science Ltd. All rights reserved. PlI:S0266-3538(99)00011-1
Fibre strength parameters measured in situ for ceramic-matrix composites tested at elevated temperature in vacuum and in air Ian J. Davies a , Takashi Ishikawa a,*, Masaki Shibuya b, Tetsuro Hirokawa c a Airframe Division, National Aerospace Laboratory, 6-13-1 Ohsawa, Mitaka-Shi, Tokyo 181, Japan bCorporate Research and Development, Ube Industries Ltd, 1978-10 Kogushi, Ube-Shi, Yamaguchi 755, Japan c Research and Development Department, Industrial Textile Division, Shikibo Ltd, 1500-5 Shibahara-Minami, Yokaichi-Shi, Shiga 527, Japan Received 3 July 1997; received in revised form 15 July 1998; accepted 7 January 1999 Abstract In situ ®bre fracture characteristics have been investigated for Si±Ti±C±O ®bres after tensile testing up to 1380C in vacuum and in air. Specimens tested in air at 1100 and 1200C generally had ¯at fracture surfaces with less than 20% of ®bres exhibiting fracture mirrors: this is attributed to oxygen ingress into the ®bre bundles. Fibre strength characteristics normalised to a 10ÿ3 m gauge length indicated that ®bres tested in air at elevated temperature have signi®cantly lower strengths and average Weibull parameter, m, compared to the room-temperature, 1200 and 1300C/vacuum cases, and this is attributed to oxygen damage of the ®bre together with oxidation of the ®bre/matrix interface. The ®bre/matrix interface shear strength, , was low for the room-temperature specimens and increased slightly with temperature when tested in vacuum, possibly as a result of a change in the thermal mismatch between ®bres and matrix. Values of for specimens tested at 1100 and 1200C in air were an order of magnitude greater than those for room-temperature specimens, indicating a signi®cant degree of oxidation damage at the ®bre/matrix interface to have occurred. # 1999 Elsevier Science Ltd. All rights reserved. 1. Introduction Ceramic-matrix composites (CMCs) that utilise continuous ®bres as reinforcement have many potential high-temperature structural applications, particularly in the area of space re-entry vehicles. Recent advances include production of 3-D woven composites based on the SiC/SiC system that possess short-term tensile strengths of nearly 400 MPa in vacuum at room temperature and 1200C with tensile strains to failure in excess of 1% [1±3]. However, the mechanical properties of CMCs often degrade at elevated temperature in the presence of oxygen which is known to attack the ®bre/ matrix interface. Although studies have shown that sealing the surface of CMCs with a glass-based compound may allow similar mechanical properties to be achieved at elevated temperature in air and in vacuum [4], improvement of the ®bre/matrix interface is still a major area of research. Recent work [5,6] has indicated that several important composite mechanical properties may be predicted from a knowledge only of the ®bre/matrix interface shear strength, , and ®bre properties such as the radius, r, and the Weibull strength parameters, So and m. Values for So and m may be obtained by measuring the in situ strength of ®bres and ®tting a two-parameter Weibull curve [7] to the resulting data, i.e. F 1 ÿ eÿ S So m 1 where F is the cumulative failure probability of ®bres at a stress, S, and So and m are empirical constants known as the Weibull strength parameters. Although the ex situ strength of ceramic ®bres is often well known, there is a relative dearth of data concerning in situ ®bre strength. One method of estimating in situ ®bre strength involves the measurement of mirror radiiÐa feature often present on the fracture surface of ceramic ®bres. Fig. 1 illustrates such a fracture mirror observed on the fracture surface of a Tyranno1 Si±Ti±C±O ®bre. It can be seen that the fracture mirror is a smooth region adjacent to the initiating defect in the ®bre and is surrounded by a region of multiple fracture planes. Past research has shown most fracture mirrors to initiate at ¯aws present on the surface of the ®bre [8,9] Composites Science and Technology 59 (1999) 801±811 0266-3538/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S0266-3538(99)00011-1 * Corresponding author
I.. Davies et al. Composites Science and Technology 59(1999)801-811 which is vulnerable to degradation during fibre and approximations for So and m with the correct form of composite processing. That fracture initiation does not the relationship being [5] normally occur within the fibre bulk suggests the aver- ge flaw size and or flaw density to be significantly lar- F=l ger on the fibre surface compared to that within the fibre bulk. A common equation linking mirror radius, with S" and m being the uncorrected Weibull strength Im, to fibre strength, S is of the form parameters In order to obtain the relevant in situ fibre strength S (2) parameters, So and m, it is required to use correction factors that have been determined [5 to be of the form shown in Fig. 2. It may be observed from Fig. 2 that Eqs (1)and (3)gives similar results when m A 4 but values for So and m obtained from Eq (1)for the case of m< 4 will Am= bric respectively underestimate and overestimate actual valu where Am and Bm are empirical constants and Kic is the It should be noted that Eq 3)assumes no knowledge racture toughness of the fibre. a value of Kic a 1 of the specimen gauge length. However, the value of S MPa m/ 2 has been estimated for Nicalon"SiC-based obtained depends strongly on the specimen gauge length fibres [ 8] whilst values of 3. 5 [10]and 2.51 [11] have been with large specimens having reduced strengths com- suggested for Bm. Although values of Bm and Kic have not pared to smaller specimens. It is thus necessary when been determined for Tyranno Si-Ti-C-O fibres, they comparing S, for different data sets to normalise differ might be expected to be similar to those for Nicalon" ent gauge lengths to a standard gauge length, Lo, by fibres as the microstructure and chemistry of Tyranno" using the relationship [9] and Nicalon"fibres are alike in many respects. It has recently been shown that Eq. (1) provides only where S, is the predicted value of S, at the standard gauge length, L', and Lo is the gauge length for a spe cimen with Weibull strength parameters So and m It has been suggested that Lo=10-3 m is an appro priate standard gauge length for CMCs as this is the order of the fibre pull-out length. The reason why fibre th is significant for CMCs is that fo posites that fail as a result of multiple matrix cracking as is the case for most“good”CMCs, the gradual transfer of stress from matrix to fibre away from the sam2 4 um 1.1 (b) E0.8 0.7 0.6 um Weibull modulus . m Fig. 1. Scanning electron micrographs illustrating a typical fractu Fig. 2. Relationship between Weibull scale parameters(S, mr) mirror observed on the surface of Tyranno" Si-TiC-O fibres:(a) mined from fracture mirror data and underlying fibre strength general view, and(b) detailed view of fracture mirror. meters(So, m)[5
which is vulnerable to degradation during ®bre and composite processing. That fracture initiation does not normally occur within the ®bre bulk suggests the average ¯aw size and/or ¯aw density to be signi®cantly larger on the ®bre surface compared to that within the ®bre bulk. A common equation linking mirror radius, rm, to ®bre strength, S, is of the form: S Am rm p 2 with Am BmKIC where Am and Bm are empirical constants and KIC is the fracture toughness of the ®bre. A value of KIC 1 MPa m1/2 has been estimated for Nicalon1 SiC-based ®bres [8] whilst values of 3.5 [10] and 2.51 [11] have been suggested for Bm. Although values of Bm and KIC have not been determined for Tyranno1 Si±Ti±C±O ®bres, they might be expected to be similar to those for Nicalon1 ®bres as the microstructure and chemistry of Tyranno1 and Nicalon1 ®bres are alike in many respects. It has recently been shown that Eq. (1) provides only approximations for So and m with the correct form of the relationship being [5]: F 1 ÿ eÿ S S m 3 with S and m being the uncorrected Weibull strength parameters. In order to obtain the relevant in situ ®bre strength parameters, So and m, it is required to use correction factors that have been determined [5] to be of the form shown in Fig. 2. It may be observed from Fig. 2 that Eqs. (1) and (3) gives similar results when m 4 but values for So and m obtained from Eq. (1) for the case of m < 4 will respectively underestimate and overestimate actual values. It should be noted that Eq. (3) assumes no knowledge of the specimen gauge length. However, the value of S obtained depends strongly on the specimen gauge length with large specimens having reduced strengths compared to smaller specimens. It is thus necessary when comparing So for dierent data sets to normalise dierent gauge lengths to a standard gauge length, L0 o, by using the relationship [9]: S0 o So Lo L0 o 1 m 4 where S0 o is the predicted value of So at the standard gauge length, L0 o, and Lo is the gauge length for a specimen with Weibull strength parameters So and m. It has been suggested that L0 o 10ÿ3 m is an appropriate standard gauge length for CMCs as this is the order of the ®bre pull-out length. The reason why ®bre pull-out length is signi®cant for CMCs is that, for composites that fail as a result of multiple matrix cracking (as is the case for most ``good'' CMCs), the gradual transfer of stress from matrix to ®bre away from the Fig. 1. Scanning electron micrographs illustrating a typical fracture mirror observed on the surface of Tyranno1 Si±Ti±C±O ®bres: (a) general view, and (b) detailed view of fracture mirror. Fig. 2. Relationship between Weibull scale parameters (S, m) determined from fracture mirror data and underlying ®bre strength parameters So; m [5]. 802 I.J. Davies et al. / Composites Science and Technology 59 (1999) 801±811
I.. Davies et al. Composites Science and Technology 59(1999)801-811 crack front means that the fibres are effectively tested at reason, the present paper will investigate in situ fibre a gauge length, 8c, independent of the composite speci- properties for composites based on the SiC/SiC system men gauge length and deduced to have the form [6, 12]: after tensile testing up to 1380 C in vacuum and air h) rS (m) 2. Experimental procedure with The composite investigated in this report was on the siC/Sic system with Tyranno" Si-Ti-C-O (m)≈0716+m6f0rm≥1 fibres being utilised that had been surface-modi order to improve interface properties. Ex situ strength data for standard and surface-modified LoxM fibres [14] where() is the mean fibre pull-out length after com- has been presented in Table I whilst Auger depth pro- ure files at the fibre surface [15] are shown in Fig. 3. Normalising So values for in situ fibres would be Increases in So and m following surface modification xpected to give reasonably accurate values for S, at were presumably due to a reduction in the average Lo= 10-m as Sc= Lo g Lo. However, some caution should be used when Lo > L as is often the case for ex situ fibre strength where L, may be typically 25x10-3 Table 1 m. In such cases, researchers [13] have discovered Ex situ strength parameters for TyrannoSH-Ti-C-O fibres calculated from data in Ref (14 appreciable differences(up to 50%) when comparing values of So obtained from experimental gauge lengths Tyranno Si-Ti-C-0 So(GPa) Le (le=lo) and those predicted from Eq.(4)for LoxM fibre (GPa) gauge length of Lo after testing at a gauge length LoStandard 3.32(±0.01)10.84(±0.32)4.47(±0.13) (Le》Le) Surface-modified4.23(±0.01)12.54(±0.39)5.47(±0.17) Once values of So and m have been established, they may be used to estimate the ultimate tensile strength of a composite, SUTs, that failed through multiple matrix cracking using the following equation [5] (a) 2+1 m+2 m+2 where V is the fibre volume fraction in the direction of Comparison between experimental and predicted alues for suts has met with some success, though deviations of 20-30% have been seen in other systems [9]. From Eq.(6)it may be observed that composite strength is determined essentially by the fibre archi- 100 tecture, V, fibre properties, So, and m, and also the interfacial shear strength, T, which can be derived from (b) rearrangement of Eq (5)such that [6] ri(m)So T 4(h Thus, in situ observation of fibres makes possible the derivation of s nd t that composite properties. Although the theory discussed above may be useful in correlating microscopic and macroscopic properties, only 020406080100120140160 limited data exist in situ observation of fibre properties in CMCs. In particular, only a few studies have Distance from fibre surface(nm) determined in situ fibre properties for CMCs after testing Fig 3. Auger depth profiles for Tyranno* LoxM Si-Ti-C-o fibres in different atmospheres at elevated temperature. For that (a)standard, and(b) surface-modified [15]
crack front means that the ®bres are eectively tested at a gauge length, c, independent of the composite specimen gauge length and deduced to have the form [6,12]: c 4hhi l m rSo 5 with l m 0:716 1:36 m0:6 for m51 where hhi is the mean ®bre pull-out length after composite failure. Normalising So values for in situ ®bres would be expected to give reasonably accurate values for S 0 o at L0 o 10ÿ3 m as c Lo L0 o. However, some caution should be used when Lo L0 o as is often the case for ex situ ®bre strength where Lo may be typically 2510ÿ3 m. In such cases, researchers [13] have discovered appreciable dierences (up to 50%) when comparing values of So obtained from experimental gauge lengths Le Le L0 o and those predicted from Eq. (4) for a gauge length of L0 o after testing at a gauge length Lo Le Le. Once values of So and m have been established, they may be used to estimate the ultimate tensile strength of a composite, SUTS, that failed through multiple matrix cracking using the following equation [5]: SUTS VfSo 2 m 2 1 m1 m 1 m 2 6 where Vf is the ®bre volume fraction in the direction of loading. Comparison between experimental and predicted values for SUTS has met with some success, though deviations of 20±30% have been seen in other systems [9]. From Eq. (6) it may be observed that composite strength is determined essentially by the ®bre architecture, Vf, ®bre properties, So, and m, and also the interfacial shear strength, , which can be derived from rearrangement of Eq. (5) such that [6]: rl mS0 4hhi 7 Thus, in situ observation of ®bres makes possible the derivation of So, m, and , that eectively control most composite properties. Although the theory discussed above may be useful in correlating microscopic and macroscopic properties, only limited data exist concerning in situ observation of ®bre properties in CMCs. In particular, only a few studies have determined in situ ®bre properties for CMCs after testing in dierent atmospheres at elevated temperature. For that reason, the present paper will investigate in situ ®bre properties for composites based on the SiC/SiC system after tensile testing up to 1380C in vacuum and air. 2. Experimental procedure The composite investigated in this report was based on the SiC/SiC system with Tyranno1 Si±Ti±C±O LoxM ®bres being utilised that had been surface-modi®ed in order to improve interface properties. Ex situ strength data for standard and surface-modi®ed LoxM ®bres [14] has been presented in Table 1 whilst Auger depth pro- ®les at the ®bre surface [15] are shown in Fig. 3. Increases in So and m following surface modi®cation were presumably due to a reduction in the average Table 1 Ex situ strength parameters for Tyranno1 Si±Ti±C±O ®bres calculated from data in Ref. [14] Tyranno Si±Ti±C±O LoxM ®bre So (GPa) m S 0 o L0 o 10ÿ3 m (GPa) Standard 3.32 (0.01) 10.84 (0.32) 4.47 (0.13) Surface-modi®ed 4.23 (0.01) 12.54 (0.39) 5.47 (0.17) Fig. 3. Auger depth pro®les for Tyranno1 LoxM Si±Ti±C±O ®bres: (a) standard, and (b) surface-modi®ed [15]. I.J. Davies et al. / Composites Science and Technology 59 (1999) 801±811 803
I.. Davies et al. Composites Science and Technology 59(1999)801-811 defect size at the fibre surface caused by the surface tested specimens were similar at room temperature and reatment. The chemical depth profile in Fig 3(b)shows 1200oC(400 MPa), followed by a gradual decrease of the surface-modified LoxM fibre to possess a 10 nm approximately 50% until 1380oC. Tensile strain to fail- SiOx-rich layer at the surface surrounding an inner 40 ure was approximately 1. 2%for specimens tested at nm carbon-rich layer. The justification for such a sur- room temperature and up to 1380.C in vacuum face chemistry is that the outer SiOx- rich layer will bond Although a suitable strain measurement technique for strongly to the matrix with the 40 nm carbon layer specimens tested in air at elevated temperature was not effectively acting as the fibre/matrix interface. In this available, values of typically <0.05% would be expec- case, the fibres would be expected to fail at the carbon ted when considering the reduced tensile strength of layer within the fibre surface [16] rather than at the specimens tested in air at elevated temperature(Fig 3) actual fibre/matrix interface that is usually the case in However, it should be emphasised that specimens inves- CMCS tigated in this report possessed no oxidation protection Prior to matrix densification the fibres were woven system. Tensile tests at elevated temperature in air for nto an orthogonal 3-D structure with fibre volume similar specimens, but surface sealed using a proprietary fractions in the x, y, and z directions being 0.19, 0.19, glass-based technique, indicate tensile strength to be simi- and 0.02, respectively. Weaving technology was utilised lar to that for unsealed specimens tested in vacuum [4] for this composite as it possesses great versatility as Following tensile failure, specimen fracture surfaces regards shape and dimension control [17], which should were investigated using a JEOL JSM-6300F scanning reduce machining costs in final applications electron microscope(SEM). The general nature of the Matrix densification consisted of a polymer similar to composite fracture surface was assessed whilst a detailed polytitanocarbosilane(PTCS) that was impregnated study of fibre pull-out behaviour is reported elsewhere into the fibre preform and pyrolysed to form a matrix [19]. Fibre fracture surfaces were characterised, with the similar in chemistry to that of the fibres. Eight cycles of fracture mode and flaw mirror radius being noted impregnation and pyrolysis were required to maximise Between 100 and 800 fibres were examined for each test the composite density [16] condition whilst in situ fibre properties were derived Tensile testing was undertaken with the specimen with the aid of Eqs.(2H(7) axis parallel to the loading direction at temperatures between room temperature and 1380 C in vacuum and from room temperature to 1200 C in air For specimens 3. Results and discussion tested at elevated temperature, heating rates between 300C and the specified test temperature were approxi- 3.I. Microstructural observation mately 0. 75C s-I whilst failed specimens were furnace cooled at an estimated rate above 1000c of 3. 3C 3. .1. Room temperature and 1200 C in vacuum The total time spent at the test temperature was believed The majority of fibres within specimens tested at to be approximately 600 s-further experimental details room temperature and 1200C in vacuum possessed being given elsewhere [15, 18 fine-grained structures with a well-defined mirror zone The tensile strengths [2] of specimens examined in this and crackle region that originated at the fibre surface report are presented in Fig 4. The strengths of vacuum-(as indicated in Fig. 1). Surface flaws thus controlled fibre strength under these conditions with relatively few fibres failing as a result of internal flaws for the room temperature case. The percentage of fibres failing at internal flay her for the 1200@C case compared to room temperature as indicated in Table 2. However, this appeared to have no effect on composite tensile ￡30 strength shown in Fig. 4, which would be consistent with fibres failing at the surface and in the bulk having similar strength distributions; this will be the topic of ■ vacuum further research. An example of a fibre that failed 100 through an internal flaw during testing at 1200oC in vacuum is presented in Fig. 5 with a mirror zone also being present around the flaw. From Table 2 it can be 10001100120013001400 seen that the failure mode could not be determined for Test temperature(C) 13% of fibres that failed at room temperature Although these fibres had no fracture mirrors visible, it Fig. 4. Tensile strength [2] of SiC/SiC-based specimens tested in was concluded that they probably did fail due to surface vacuum and air up to 1380.C. Note the non linear temperature scale flaws and may have represented the strongest population
defect size at the ®bre surface caused by the surface treatment. The chemical depth pro®le in Fig. 3(b) shows the surface-modi®ed LoxM ®bre to possess a 10 nm SiOx-rich layer at the surface surrounding an inner 40 nm carbon-rich layer. The justi®cation for such a surface chemistry is that the outer SiOx-rich layer will bond strongly to the matrix with the 40 nm carbon layer eectively acting as the ®bre/matrix interface. In this case, the ®bres would be expected to fail at the carbon layer within the ®bre surface [16] rather than at the actual ®bre/matrix interface that is usually the case in CMCs. Prior to matrix densi®cation the ®bres were woven into an orthogonal 3-D structure with ®bre volume fractions in the x, y, and z directions being 0.19, 0.19, and 0.02, respectively. Weaving technology was utilised for this composite as it possesses great versatility as regards shape and dimension control [17], which should reduce machining costs in ®nal applications. Matrix densi®cation consisted of a polymer similar to polytitanocarbosilane (PTCS) that was impregnated into the ®bre preform and pyrolysed to form a matrix similar in chemistry to that of the ®bres. Eight cycles of impregnation and pyrolysis were required to maximise the composite density [16]. Tensile testing was undertaken with the specimen y-axis parallel to the loading direction at temperatures between room temperature and 1380C in vacuum and from room temperature to 1200C in air. For specimens tested at elevated temperature, heating rates between 300C and the speci®ed test temperature were approximately 0.75C sÿ1 whilst failed specimens were furnacecooled at an estimated rate above 1000C of 3.3C sÿ1 . The total time spent at the test temperature was believed to be approximately 600 sÐfurther experimental details being given elsewhere [15,18]. The tensile strengths [2] of specimens examined in this report are presented in Fig. 4. The strengths of vacuumtested specimens were similar at room temperature and 1200C (400 MPa), followed by a gradual decrease of approximately 50% until 1380C. Tensile strain to failure was approximately 1.2% for specimens tested at room temperature and up to 1380C in vacuum. Although a suitable strain measurement technique for specimens tested in air at elevated temperature was not available, values of typically <0.05% would be expected when considering the reduced tensile strength of specimens tested in air at elevated temperature (Fig. 3). However, it should be emphasised that specimens investigated in this report possessed no oxidation protection system. Tensile tests at elevated temperature in air for similar specimens, but surface sealed using a proprietary glass-based technique, indicate tensile strength to be similar to that for unsealed specimens tested in vacuum [4]. Following tensile failure, specimen fracture surfaces were investigated using a JEOL JSM-6300F scanning electron microscope (SEM). The general nature of the composite fracture surface was assessed whilst a detailed study of ®bre pull-out behaviour is reported elsewhere [19]. Fibre fracture surfaces were characterised, with the fracture mode and ¯aw mirror radius being noted. Between 100 and 800 ®bres were examined for each test condition whilst in situ ®bre properties were derived with the aid of Eqs. (2)±(7). 3. Results and discussion 3.1. Microstructural observation 3.1.1. Room temperature and 1200C in vacuum The majority of ®bres within specimens tested at room temperature and 1200C in vacuum possessed ®ne-grained structures with a well-de®ned mirror zone and crackle region that originated at the ®bre surface (as indicated in Fig. 1). Surface ¯aws thus controlled ®bre strength under these conditions with relatively few ®bres failing as a result of internal ¯aws for the room temperature case. The percentage of ®bres failing at internal ¯aws was higher for the 1200C case compared to room temperature as indicated in Table 2. However, this appeared to have no eect on composite tensile strength shown in Fig. 4, which would be consistent with ®bres failing at the surface and in the bulk having similar strength distributions; this will be the topic of further research. An example of a ®bre that failed through an internal ¯aw during testing at 1200C in vacuum is presented in Fig. 5 with a mirror zone also being present around the ¯aw. From Table 2 it can be seen that the failure mode could not be determined for 13% of ®bres that failed at room temperature. Although these ®bres had no fracture mirrors visible, it was concluded that they probably did fail due to surface ¯aws and may have represented the strongest population Fig. 4. Tensile strength [2] of SiC/SiC-based specimens tested in vacuum and air up to 1380C. Note the non-linear temperature scale. 804 I.J. Davies et al. / Composites Science and Technology 59 (1999) 801±811
I.. Davies et al. Composites Science and Technolog y 59(1999)801-811 Table 2 In situ fibre fracture surface characteristics for Tyranno* Si-THC-O fibres Test condition Fracture mirror Fracture mirror Flat fracture Undetermined Total (surface)%(N (internal)%(M surface %( %(N 75(114) 13(19) 100(151) 200°C/ vacuum 10(11) l(1 00(112) 1300° C/vacuum 3(5) 1200°C/air 100(788) a 3 um 4 um (b) (b) lm 2 um Fig. 5. Scanning electron micrographs illustrating the fracture surface Tyranno"ShTHC-O fibre tested in situ at 1200C in vacuum tha Fig. 6. Scanning electron micrographs illustrating the fracture surface failed due to an internal flaw:(a) general view, and (b) detailed view of f a Tyranno"Si-Ti-C-O fibre tested in situ at 1300.C in vacuum that fracture mirror failed due to a surface flaw:(a) general view, and (h) detailed view of fracture mirror of fibres that"shattered"upon failure. Likewise, fibres evolution of Co, with associated decreased tensile characterised as"flat fracture surface"in Table 2 for strength [21]. That this phenomenon was not observed the room temperature test condition possessed smooth in current specimens until 1300C was believed, in part features that suggested them to be the weakest group of to be due to the relatively short time above 1000C during testing compared to previous resea observed in Fig. 6 is a fracture mirror that was typical 3..2.l300° c in vacu of those seen in 84% of fibres at 1300C in vacuum Whereas fibres tested at 1200'C in vacuum indicated (Table 2). The number of fibres that failed due to inter- no obvious grain growth compared to room tempera- nal flaws was about 12% and only slightly larger than ture fibres, those tested at 1300@C exhibited significant that at 1200C (10%)indicating that the voids observed grain growth and finely distributed voids(Fig. 6). Such in Fig. 6 were not large enough to significantly further a phenomenon is known to occur in SiC-based fibres challenge the surface flaw-induced failure mode follow- held in an inert atmosphere above 1000.C[20] and ing the initial increase between room temperature and attributed to chemical decomposition of the fibre and 1200oC
of ®bres that ``shattered'' upon failure. Likewise, ®bres characterised as ``¯at fracture surface'' in Table 2 for the room temperature test condition possessed smooth features that suggested them to be the weakest group of ®bres. 3.1.2. 1300C in vacuum Whereas ®bres tested at 1200C in vacuum indicated no obvious grain growth compared to room temperature ®bres, those tested at 1300C exhibited signi®cant grain growth and ®nely distributed voids (Fig. 6). Such a phenomenon is known to occur in SiC-based ®bres held in an inert atmosphere above 1000C [20] and attributed to chemical decomposition of the ®bre and evolution of CO, with associated decreased tensile strength [21]. That this phenomenon was not observed in current specimens until 1300C was believed, in part, to be due to the relatively short time above 1000C during testing compared to previous researchers. Also observed in Fig. 6 is a fracture mirror that was typical of those seen in 84% of ®bres at 1300C in vacuum (Table 2). The number of ®bres that failed due to internal ¯aws was about 12% and only slightly larger than that at 1200C (10%) indicating that the voids observed in Fig. 6 were not large enough to signi®cantly further challenge the surface ¯aw-induced failure mode following the initial increase between room temperature and 1200C. Table 2 In situ ®bre fracture surface characteristics for Tyranno1 Si±Ti±C±O ®bres Test condition Fracture mirror (surface) % (N) Fracture mirror (internal) % (N) Flat fracture surface % (N) Undetermined % (N) Total % (N) Room temperature 75 (114) 1 (2) 11 (16) 13 (19) 100 (151) 1200C/vacuum 89 (100) 10 (11) 0 (0) 1 (1) 100 (112) 1300C/vacuum 84 (132) 12 (19) 1 (1) 3 (5) 100 (157) 1100C/air 18 (128) 0 (0) 82 (570) 0 (0) 100 (698) 1200C/air 18 (64) 0 (0) 92 (724) 0 (0) 100 (788) Fig. 6. Scanning electron micrographs illustrating the fracture surface of a Tyranno1 Si±Ti±C±O ®bre tested in situ at 1300C in vacuum that failed due to a surface ¯aw: (a) general view, and (h) detailed view of fracture mirror. Fig. 5. Scanning electron micrographs illustrating the fracture surface of a Tyranno1 Si±Ti±C±O ®bre tested in situ at 1200C in vacuum that failed due to an internal ¯aw: (a) general view, and (b) detailed view of fracture mirror. I.J. Davies et al. / Composites Science and Technology 59 (1999) 801±811 805