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SiC composites with a BN interphase
Acta mater. 48 (2000) 4609–4618 www.elsevier.com/locate/actamat
THE CONCEPT OF A STRONG INTERFACE APPLIED TO SiC/SiC COMPOSITES WITH A BN INTERPHASE F. REBILLAT, J. LAMON† and A. GUETTE Laboratoire des Composites Thermostructuraux, UMR 5801, CNRS-SNECMA-CEA-UB1, 3, alle´e de la Boe´tie, 33600 Pessac, France
Abstract—Strong interfaces have been shown to allow improvement of the mechanical properties of ceramic matrix composites (CMC). The concept of a strong interface has been established in SiC/SiC composites with pyrocarbon (PyC) or multilayered (PyC/SiC) fiber coatings (also referred to as interphases). The present paper reports an attempt directed at applying the concept of a strong interface to SiC/SiC composites with a BN coating (referred to as SiC/BN/SiC). Fiber bonding and frictional sliding were investigated by means of push-out tests performed on 2D-composites as well as on microcomposite samples, and tensile tests performed on microcomposites. The stress–strain behavior of the SiC/BN/SiC composites and microcomposites is discussed with respect to interface characteristics and location of debonding either in the coating or in the fiber/coating interface. 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Interphase; Interface; Composites
1. INTRODUCTION
The fiber–matrix interfacial domain is a critical part of composites because load transfers from the matrix to the fiber and vice versa must occur through the interface. Therefore, it exerts a profound influence upon the mechanical behavior and the lifetime. Thus, it may be expected that composites could be tailored as a function of end use applications through optimization of interfaces. In fiber-reinforced ceramic composites, most authors promote the concept of weak interfaces to increase fracture toughness. The major contribution to toughness is attributed to crack bridging and fiber pull-out [1, 2]. Weak interfaces are detrimental to composite strength. A high strength requires efficient load transfers which are obtained with strong interfaces. This implies short debond cracks and/or significant sliding friction. These latter requirements, to be met for strong composites, are therefore incompatible with the former ones for tough composites, if toughening is based solely upon the above mentioned weak interface-based mechanisms. Fiber/matrix bonding results from diffusion or chemical reactions (chemical bonding) or from fiber clamping by residual stresses induced by thermal
† To whom all correspondence should be addressed. Tel.: ⫹33-556-844-703; fax: ⫹33-556-841-225. E-mail address: [email protected] (J. Lamon)
expansion mismatch. Fiber/matrix interfaces in the most advanced ceramic matrix composites consist of a thin coating layer (less than 1 µm thick) of one or several materials deposited on the fiber (interphase). Recently, SiC/SiC composites with strong interfaces have been developed. The coating/fiber bond was significantly stronger when fibers had been previously treated [3–5]. Features of the mechanical behavior of SiC/SiC composites with strong fiber/coating interfaces has been examined in several papers [4, 6–10]. Experiments as well as models have demonstrated that a strong interface is beneficial to the strength, the toughness, the lifetime and the creep resistance [4, 6– 11]. By contrast, weak interfaces are shown to be detrimental. The concept of strong interfaces has been established on SiC/C/SiC composites with PyC and multilayered (PyC/SiC) fiber coatings. In the present paper, it is applied to SiC/BN/SiC composites with boron nitride fiber coatings. BN is foreseen to be an alternative fiber coating to improve the oxidation resistance of ceramic matrix composites at high temperature.
2. FEATURES OF STRONG INTERFACES VS WEAK INTERFACES
In order to properly introduce the concept of a strong interface, we recall first the basic features of interface phenomena in CMCs subject to an essentially tensile load. These phenomena influence the
1359-6454/00/$20.00 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 0 ) 0 0 2 4 7 - 0
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mechanical response of composites described by the stress–strain curve. Fiber debonding results from the deflection of the cracks that initiate in the matrix (Fig. 1). Then sliding of the fiber debonded in the interface determines the load transfers from the fiber to the matrix and vice versa. The fiber sliding is influenced by the misfit strain, the associated radial component of the thermally induced residual stress-field, surface roughness and debond length. Weak interfaces debond easily. A single long debond crack is located at the surface of the fibers in those composites exhibiting weak interfaces (adhesive failure type, Fig. 1). As a consequence of small interface shear stresses, load transfers through the debonded interfaces are poor. The matrix becomes subjected to lower stresses and the volume of matrix that may experience further cracking is reduced by the presence of long debonds. The cracks are generally widely opened, whereas the crack spacing distance at saturation as well as the pull out length tend to be rather long (>100 µm). Toughening results essentially from sliding friction along the debonds. However, due to poor load transfers and long debonds, the fibers carry most of the load, which reduces the composite strength. The corresponding tensile stress–strain curve exhibits a short curved domain limited by a stress at matrix cracking saturation which is significantly smaller than ultimate strength (Fig. 2). In the presence of stronger fiber/coating bonds, the matrix cracks are deflected within the coating
Fig. 2. Typical tensile stress–strain behaviors measured on 2D SiC/SiC composites possessing PyC based interphases and fabricated from untreated or treated Nicalon (ceramic grade) fibers: (a) strong fiber/coating interfaces and (b) weak fiber/coating interfaces.
(cohesive failure type, Fig. 1), into short and branched multiple cracks [4, 12]. Short debonds as well as improved load transfers allow further cracking of the matrix via a scale effect [6, 7] leading to a higher density of matrix cracks (which are slightly opened). Sliding friction within the coating as well as multiple cracking of the matrix increase energy absorption, leading to toughening. Limited debonding and improved load transfers reduce the load carried by the fibers, leading to strengthening. The associated tensile stress strain curve exhibits a wide curved domain and the stress at matrix cracking saturation is close to ultimate failure (Fig. 2). Table 1 gives various values of the interfacial shear stresses measured using various methods on SiC/SiC composites with PyC-based fiber coating. It can be noticed that the interfacial shear stresses range between 10 and 20 MPa for the weak interfaces whereas they are larger than 100–300 MPa for the strong interfaces. Additional data can be found in [4, 7, 8, 24, 29]. 3. SiC/BN/SiC COMPOSITES: TESTING METHODOLOGY AND MICROSTRUCTURAL ANALYSES
3.1. Specimen preparation
Fig. 1. Schematic diagram showing crack deflection when the fiber coating/interface is (a) strong or (b) weak.
SiC/BN/SiC microcomposites and woven composites were manufactured via chemical vapor infiltration [13]. They were reinforced with either asreceived or treated (proprietary treatment, SNECMA/SEP, Bordeaux) SiC Nicalon fibers (NL 202 grade). The SiC/BN/SiC microcomposites consist of a single fiber (15 µm diameter), coated with a boron nitride layer (0.3–0.9 µm thick) and a SiC matrix deposited by CVD. They are representative of their counterparts in the 2D woven composites, since they are produced using identical chemical vapor deposition conditions [13]. A single or a bilayered BN fiber coating was deposited from a BF3, NH3, Ar gas mixture (Table
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Table 1. Interfacial shear stress (MPa) measured using various methods on 2D-SiC/SiC composites with PyC based fiber coatings and reinforced with either as-received or treated fibers SiC/C/SiC SiC/(C/SiC)n/SiC
Untreated fibers 2D woven Microcomposites Minicomposites 2D woven
Treated fibers 2D woven
Interphase
Crack spacing [30]
Crack spacing [31]
Tensile tests (hysteresis loops) [24]
PyC (0.1) PyC(0.1) PyC(0.1) PyC(0.5) (PyC/SiC)2 (PyC/SiC)4
12
8
0.7 3 21–115 4 2 9
PyC(0.1) PyC(0.5) (PyC/SiC)2 (PyC/SiC)4
203
140
190 370 150 90
Tensile tests (curved domain) [7, 24]
Push-out tests (curved domain) [8, 29]
Push-out tests (plateau) [8, 29]
14–16 31 28
12–10 19.3 12.5
4–20 40–80
165–273 100–105 133 90
2). The selected processing conditions have been shown to cause minimum damage to the fibers and to improve adhesion of the BN coating onto the fibers, and the microstructure [13]. In the bi-layered coating (referred to as BN4, Table 2 and Fig. 3), the first sublayer on the fiber is BN2 type (poorly crystallized), whereas the second one is BN1 type (highly crystallized). Processing of BN2 involved the less aggressive gaseous phase, which led to a better contact between the fiber and the coating. The processing conditions of BN1 were found to be aggressive against the fibers [13]. 3.2. Push-out tests Fiber bonding and frictional sliding in the 2DSiC/BN/SiC composites were investigated by means of single fiber push-out tests [14–17]. 500 µm thick wedges were prepared using standard metallographic techniques. An interfacial test system (designed by ONERA, France) was used. The load was applied to the top of the fiber using a flat bottom diamond cone (at a constant displacement rate of 0.1 µm/s). The interface characteristics were extracted from the experimental stress–fiber end displacement curves by fitting the push-out model of Hsueh [15], as discussed in a previous paper [16]. Only a few push-out experiments could be carried out on the microcomposites owing to the difficulties involved in microcomposite handling, preparation and testing. Parallel-faced strips were cut out of the microcomposites which had been previously embedded in glass [18] (microcomposites 2) or in a ceramic cement [19] (microcomposites 2 and 4). A first series
Fig. 3. Images of BN coatings: (a) TEM-image and DEASpicture of the BN1 coating showing the three-dimensional ordered hexagonal structure; (b) SEM-image of bi-layered BN interphase (BN4) in a SiC/BN/SiC microcomposite showing poorly ordered BN2 layer and the BN1 layer with a threedimensional ordered hexagonal structure.
Table 2. Main characteristics of the BN fiber coatings [13] Batch
1 2 4
BN coating Number of Degree of BN layers crystallization BN1 BN2 BN4
1 1 2
High Low BN1⫹BN2
Coating adhesion Strong Weak Strong
of push out tests on microcomposites 2 used a Vickers diamond probe [18]. A correction for indentor displacement was done. Most of the tests were pushin rather than push-out tests. Push-out could not be achieved on the specimens with a thickness exceeding 200 µm. A number of push-in curves was unusable
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for analysis in that they exhibited features that were inconsistent with the model. A second series of pushin tests on microcomposites 2 was then performed on thicker samples (290 µm) using a flat-bottomed cone. Push-out tests were successfull on samples of microcomposites 4 (thickness ⬇190 µm). The interfaces characteristics were extracted by fitting the Hsueh’s model [15] to the push-in curves or to the curved domain and to the plateau of the push-out curves. 3.3. Tensile tests Five microcomposites per batch were tested in tension by using a specific table-model testing machine designed and developed for fiber testing [20]. The single fiber tensile test procedure based on window frames with appropriate gauge lengths (generally 10 mm) was employed [21, 22]. The microcomposites were loaded up to failure, either monotonically or with unloading–reloading cycles at a low strain rate (0.1% mn⫺1). The interface characteristics including the shear stress (t), the debond energy (Gic) and the debond length (ld) were extracted from the stress–strain curves [22–24] and from hysteresis loops on unloading–reloading [23, 24]. Independent models were used in order to assess the results. These models are referred to as LRLC, CLR and LRE according to authors’ names [22–24]. They derive from modelling the tensile load–displacement behavior (LRLC model) or the hysteretic behavior (CLR and LRE models) of microcomposites: the CLR model determines the energy dissipated in the friction phenomena whereas the LRE one determines the crack opening displacement during unloading–reloading cycles. After ultimate failure, the microcomposites were examined using scanning electron microscopy (SEM). The composition of the surface of fibers was determined from Auger electron spectroscopy (AES) depthprofile analyses of the pulled-out fibers. The tensile tests on the SiC/BN/SiC woven composites (three test specimens per batch) were performed at a constant strain rate of 0.05% min⫺1. Deformations were measured using an extensometer (gauge length 25 mm). The dimensions of the test specimens were as follows: thickness 3 mm, width 8 mm, length 100 mm. 4. RESULTS
4.1. Tensile tests on the SiC/BN/SiC microcomposites 4.1.1. Stress–strain curves. Most of the tensile stress–strain curves (Fig. 4) exhibited a curved domain over a wide range of deformations (0.2– 0.9%), and rather large strains-to-failure up to 1.2% (Table 3). However, most of the microcomposites with treated fibers essentially experienced premature failures. A plateau-like behavior was observed for
Fig. 4. Tensile stress–strain curves measured on the SiC/BN/SiC microcomposites reinforced with: (a) as-received fibers and (b) treated fibers.
microcomposites 1 and 4 reinforced with as-received fibers, suggesting the presence of rather weak fiber/matrix interactions, short debonds and small densities of matrix cracks at saturation when comparing with the microcomposites 2 which exhibit a widely curved stress–strain behavior up to ultimate failure and larger stresses. Saturation of matrix cracking generally occurred at rather large deformations (>0.6%). 4.1.2. SEM fractography of microcomposites. The numbers of matrix cracks identified on the microcomposites after ultimate failure were generally comparable with the numbers of load drops or of slope decreases on the force–displacement curves (Table 3). The higher density of matrix cracks was observed in microcomposites 2 (Table 3). In the microcomposites reinforced with untreated fibers debonding was observed mainly at the fiber/BN interface. The free surface of untreated fibers, on which the BN interphase was deposited, is at least partly made of silica [3]. The resulting fiber/BN interface has been reported to correspond to a very weak bond [3, 13, 25, 26]. Similar features have been observed on SiC/C/SiC composites with a fiber coating of anisotropic PyC [5]. In the microcomposites reinforced with treated fibers, debonding was detected in the BN coating only in microcomposites 4 with a bi-layered BN coating, which indicates that the weakest link is now located in the interface between the BN sublayers. As
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Table 3. Main features of the stress–strain curves for the SiC/BN/SiC microcomposites and 2D woven composites
Interphase
Microcomposites
Composites
a b
BN1 BN2 BN4 BN1 BN2 BN4
Interphase thickness (µm)
0.28 0.27 0.29 0.5 0.3 0.5
Vf
0.42 0.76 0.47 0.40 0.40 0.40
Failure stress (MPa)
Failure strain (%)
Sa
Ta
S
792 1368 970 220 210 200
680 1813 670 32 110 210
0.55 0.85 0.99 0.58 0.38 0.5
T 0.18 1.27 0.2 0.06 0.11 0.064
Number of cracks at saturationb S
T
9 55 14
2 46 1
S⫽as-received fibers, T⫽treated fibers. Determined by SEM.
previously reported for Pyrocarbon fiber coatings [4, 5], treated fibers seem to give stronger fiber/BN bonds. However, the microcomposites with a single layer BN coating appear to be an exception to this rule, since the interface crack was detected at the fiber/BN interface. This was attributed to the presence of a weakly bonded sublayer of carbon that formed on the fibers [27]. 4.1.3. Auger electron spectroscopy analyses. AES depth-profile analyses of the pulled out fibers in microcomposites reinforced with untreated fibers, showed that the fiber surface is rich in free carbon. A layer enriched in carbon and oxygen (probably consisting of silica) is present under this carbon layer. Such a complex interfacial sequence has been already observed in 2D SiC/BN/SiC composites [26, 28]. The very thin carbon layer results from the attack of the fiber surface during BN processing [27]. 4.1.4. Extraction of interfacial properties from the stress–strain curve. The models provide comparable estimates of interfacial shear stresses for the microcomposites reinforced with untreated fibers (Fig. 5). A certain discrepancy may be observed for microcomposites 2. The interfacial shear stresses can be grouped into two distinct families (Fig. 5): t⬇5 MPa for microcomposites 4, tⱖ10 MPa for microcomposites 1 and 2, The debond energy estimates range between 1 and 8 J/m2 (Fig. 5). The interfacial characteristics determined for the microcomposites reinforced with treated fibers are shown on Fig. 6. The interfacial shear stresses obtained for microcomposites 1 and 2 are larger than those measured for the microcomposites reinforced with as-received fibers. A certain discrepancy is observed on the data extracted using the LRE model [24]: t ⫽ 400 MPa and Gic ⫽ 70 J/m2 seem to be overestimations although microcomposites 1 experienced a premature failure. The characteristics provided by the other models seem to be more realistic: 10⬍t⬍50 MPa, 0⬍Gic⬍7 J/m2 (Fig. 6).
Fig. 5. Interfacial characteristics estimated using various models for various BN interphases in SiC/BN/SiC microcomposites reinforced with untreated fibers.
4.2. Tensile tests on the 2D SiC/BN/SiC composites The stress–strain curves of the 2D SiC/BN/SiC composites also display a curved domain (Fig. 7). The strains-to-failure are smaller than those measured on the microcomposites (Table 3). They are close to 0.6% for composites 1 (reinforced with untreated fibers) and 4 (reinforced with as-received or treated fibers), whereas the other composites failed at deformations ⬍0.2%. 4.3. Push-out tests on the 2D SiC/BN/SiC composites 4.3.1. Composites reinforced with as-received fibers. The stresses to initiate and propagate the
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Fig. 6. Interfacial characteristics estimated using various models for various BN interphases in SiC/BN/SiC microcomposites reinforced with treated fibers.
debond crack as well as the interfacial shear stresses are much higher than those measured on 2D woven SiC/C/SiC composites with a PyC-based fiber coating and as received fibers [8, 29] (Table 4 and Fig. 8). As previously with the microcomposites, the largest t were obtained for composites 1 and 2. The debond stresses and the interfacial shear stresses thus appear to be insensitive to the conditions of BN processing. However, the effect of friction seems to be the most efficient in the composites 2, as indicated by the comparison of the respective values of the following specific parameters (Table 4): the applied maximum stress, the fiber-end displacement and the roughness amplitude. The smallest interfacial shear stress and roughness amplitude (A) were estimated for the bilayered BN coating (Table 4). SEM revealed the following interesting features: (i) in composites 1, the fiber surface was rather rough and the matrix surface showed a lot of pores, (ii) in composites 2, the slid surface was very smooth and debonding seemed to have occurred at the fiber/coating interface (as is usually observed in such composites [18, 26]), (iii) in composites 4, the debond crack was located in the interface between the BN sublayers (Fig. 9). 4.3.2. Composites reinforced with treated fibers. A large debond stress was estimated for composites 2 (sd ⫽ 2000 MPa, Table 4). It is much larger than
Fig. 7. Tensile stress–strain curves measured on the SiC/BN/SiC 2D composites reinforced with: (a) as-received fibers and (b) treated fibers.
those estimated for the composites reinforced with untreated fibers. Unfortunately, interfacial properties could not be extracted from the curved domain of the push-out curve, owing to bending of the sample as a result of the high load required to cause debonding. The interfacial shear stress that was extracted from the plateau is large (t⬇140 MPa). Smaller interfacial shear and debond stresses were obtained for composites 4 (Table 4). These interfacial shear stresses are close to those measured on the composites reinforced with untreated fibers. The magnitudes of debond stresses may be related to the respective stress–strain curves displayed by composites 2 and 4. The premature ultimate failure of composites 2 may correspond to the large sd value. The stress–strain behavior of composites 4, which is close to that observed for the composites reinforced with untreated fibers, may reflect the similarity in the respective interface characteristics t and sd. SEM revealed that debonding took place at the fiber/BN interface in composites 2. The contact between the coating and the fiber was not continuous, probably, as a result of the attack of the fiber by the gaseous phase. In composites 4, debonding occurred in the interface between the BN sublayers. Figure 9 shows that a BN sublayer remains bonded to the fiber. SEM examination of the fracture surface of 2D specimens tested in tension also showed a BN sublayer stuck on the fibers.
a
200 (10) 90 (15) 160 (50) 60 (15) 120 (10) 280 (10) 190 (10)
BN1 (Sa) BN2 (S) BN4 (S) BN2 (Ta) BN4 (T) BN2 (S) push-in BN4 (S) push-out
S⫽as-received fibers, T⫽treated fibers.
Microcomposites
Composites
Sample thickness (µm)
Interphase
1130 (200) 1250 (170) 1200 (330) 2000 (440) 740 (100) 1600 (400) 1100 (500)
sd (MPa) 3680 (1315) 3770 (530) 2720 (760) 2770 (350) 2670 (300) 5500 (900) 2500 (250)
smax (MPa)
0.9 (0.1) 1.6 (0.4) 0.6 (0.3)
1.35 (0.3) 1 (0.2) 1.1 (0.6)
Displacement (µm)
48 (9)
66 (9) 95 (20) 48 (9) 140 (26) 77 (9)
t (plateau) (MPa)
66 (5) 100 (40) 40 (15)
87 (22) 83 (15) 41 (13)
t (MPa)
0.06 (0.01) 0.11 (0.02) 0.04 (0.005)
0.11 (0.02) 0.09 (0.03) 0.09 (0.03)
m
68 (7) 50 (20) 45 (23)
32 (14) 45 (20) 16 (10)
A (nm)
120 (8) 110 (35) 110 (30)
120 (20) 95 (9) 105 (15)
ld (µm)
1.1 (0.4) 16 (8) 0.5 (0.4)
2.1 (1.1) 2.8 (1.4) 1.5 (1.9)
Gic (J/m2)
Table 4. Interfacial characteristics extracted from the push-out curves for the 2D-SiC/BN/SiC composites reinforced with as-received or treated fibers, and from push-in or push-out curves for the SiC/BN/SiC microcomposites reinforced with as-received fibers (standard deviation is given in brackets)
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Fig. 8. Stress–fiber end displacement single fiber push-out curves measured on a SiC/BN/SiC microcomposite and a 2D SiC/BN/SiC composite (interphase BN4) reinforced with asreceived fibers.
The debond energy estimates fall within the range obtained from the tensile tests (Fig. 5). The debond stress in microcomposites 2 is quite high (sd ⫽ 1500 MPa), as well as the interfacial shear stress (t ⫽ 90 MPa), which indicates a high resistance to debonding. This may be related to the fact that the fiber could not be pushed out the matrix, even with thin samples. For microcomposites 4, a very good agreement can be noticed between the interfacial shear stresses extracted from respectively the curved domain and the plateau of the push out curves. As previously, the lower interfacial shear stress is obtained for this microcomposite with a double layer coating. Furthermore, push out of the fiber did not require a load as high as that (smax) applied during fiber pushing in the microcomposites 2. SEM showed that debonding occurred essentially at the fiber/coating interface. Generally, really smooth slid surfaces were observed. A certain roughness was detected. 5. DISCUSSION
Fig. 9. SEM micrographs showing protruding fibers after single fiber push-out tests performed on 2D-SiC/BN/SiC composites with an interphase made of two layers of boron nitride (BN4): (a) composite reinforced with as-received fibers and (b) composite reinforced with treated fibers.
4.4. Push-out and push-in tests on microcomposites The few push-out and push-in curves that were suitable for analysis exhibited the features usually observed with 2D composites (Fig. 8). The interface parameters extracted from these curves are summarized in Table 4. The interfacial shear stresses are larger than those extracted from the tensile stress–strain curves (Fig. 5). They are in excellent agreement with those estimated on the 2D-SiC/BN/SiC composites (Table 4).
The difficulty to perform push-out tests on thin samples becomes still worse with microcomposites. Large load transfers require very thin embedded lengths in order to permit push-out. The parallel-faced strip is not the most convenient test specimen because it is too fragile, but it remains the only usable sample geometry. During a push-in, unknown maximum debond length, damage of the fiber due to high compressive stresses, bending of the sample and diamond meeting with the matrix are factors that affect the stress–strain curves, and the validity of the extracted characteristics. Then, getting well polished thin strips with a thickness ⬍200 µm is rare. This explains why the push-out tests could be carried out only on microcomposites 4 which possessed the weakest interfaces. The less aggressive gaseous phase was used for processing the BN2 coating. This coating has been shown to adhere well to the fiber [13]. The resistance to propagation of the interface crack during the pushout tests appeared to be high, which suggests efficient load transfers under a tensile load. During the tensile tests, the composites 2 experienced a premature failure whereas the microcomposites 2 exhibited a high strain at saturation. Both features are compatible with efficient load transfers. Ultimate failure of 2D composites involves additionnal phenomena and it is affected by variability in fiber strength degradation during processing. The debond stresses and the interfacial shear stresses are larger than those obtained for composites and microcomposites 4 with a bi-layered coating, and they are comparable to those estimated for composites 1 with an interphase processed in more aggressive conditions. As previously mentioned, the propagation of interface cracks during the push-out tests was easier for composites 1 than
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for composites 2. This result is also indicated by the stress–strain tensile curves of microcomposites 1 which display a plateau-like feature when comparing with those of microcomposites 2. This suggests that the fiber BN1 bond is weaker. The plateau-like feature was not observed with the 2D composites. Furthermore, microcomposites 1 failed at a smaller deformation, which may be attributed to the degradation of the fiber during BN processing. The interface between the BN sublayers in composites 4 (with as-received or treated fibers), and in microcomposites 4 with treated fibers, was found to be comparatively weaker. The fact that the associated interfacial shear stresses are the smallest is not reflected by the stress–strain curves of the 2D composites nor of the microcomposites (premature failure). The interfacial shear stresses are not distinct enough to influence the stress–strain behavior. This conclusion agrees with predictions [6, 7]. With treated fibers, the fiber/coating bond was strengthened, as shown by the estimated interfacial shear stresses and the debond stresses. In the composites 2 reinforced with treated fibers, sd (which may be considered to be commensurate with the interface bond strength) and t were significantly increased. In the composites 4 with bi-layered fiber coatings, debonding occurred in the interface between the BN sublayers. Furthermore, the corresponding shear and debond stresses are small when comparing with those pertinent to the other SiC/BN/SiC composites for which debonding took place in the fiber/BN coating interface. It is particularly interesting to compare composites 2 and 4 for which a BN2 layer is on the fiber. It may be considered that the strength of the fiber/BN2 bond was the same in both composites. Therefore it may be concluded that the interface between the BN sublayers was weaker than the fiber/BN2 bond. The tensile behavior of the 2D SiC/BN/SiC composites may be considered to agree with the interface characteristics measured using push-out tests, since close interfacial shear stresses were obtained for the composites that exhibited a comparable stress–strain behavior, whereas the largest interfacial shear stresses were observed for the composites 2 with treated fibers that experienced a premature failure. This conclusion can be also drawn for the microcomposites reinforced with as-received fibers. The interfacial shear stresses are in a good correlation with the stress–strain behavior, except for microcomposites 1. Thus, the largest interfacial shear stresses were determined for microcomposites 2 which possessed the highest density of matrix cracks at saturation (Table 3) and gave the most pronounced curvature in the stress–strain curve. However, the thermally-induced residual stresses, resulting from the thermal expansion mismatch between the fibers and the matrix, exert a certain influence. They are increased by a low volume fraction of matrix, which may explain discrepancies in the stress–strain
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behavior. The microcomposites 2 possessed the thinnest matrix layer (Table 3). Comparable interfacial characteristics and an identical tensile stress–strain behavior were obtained for earlier SiC/BN/SiC composites with as-received Nicalon fibers but a different boron nitride [26]. A difference in ultimate failure can be noticed when comparing the microcomposites and the composites. This difference must be attributed to the mechanisms involved. As previously mentioned, the ultimate failure of 2D composites involves fiber interactions and the individual break of the weakest fibers prior to instability. The ultimate failure of microcomposites is thus dictated by the filament strength (which exhibits a wide statistical distribution), whereas that of 2D composites is determined by fiber tow strength (which exhibits a limited scatter). 6. CONCLUSION
In the SiC/BN/SiC microcomposites reinforced with untreated fibers, debonding occurred between the fiber free surface and the coating. The fiber/BN interface is the weakest bond in the interfacial sequence. The presence of a very thin layer of carbon between a SiO2–C mixture layer and the BN coating was detected by AES depth profile analyses. This sublayer seems to affect the interface bond. When the BN coating was deposited on treated fibers, the fiber/coating interface was stronger, but crack deflection did not occur within the BN coating. Only bi-layered BN coatings with a weak interface between the sublayers experienced crack deviation within the coating. The interface between the BN sublayers becomes the weakest link in the interfacial sequence between the fiber and the matrix. The micromechanics based models used for the analysis of the stress–strain behavior of the microcomposites provided comparable interfacial shear stresses, despite a certain scatter in some cases. A certain discrepancy was observed for the interface fracture energy estimates, which fell within a range of small values (1–10 J/m2). Although there might be some uncertainly in the t data estimated from tensile tests on microcomposites, they indicated stronger fiber/matrix bonds in the microcomposites fabricated with treated fibers. The data determined using the Hsueh’s pushout model were about one order of magnitude larger than those extracted from the tensile stress–strain behavior of microcomposites. The origin of this discrepancy has not been elucidated in the paper. The tensile stress/strain curves determined with microcomposites were similar to those obtained for the 2D composites. Tensile tests on microcomposites may be regarded as an interesting approach to interface design. Despite the difficulties to perform the push-out tests on microcomposites, this sample geometry seems to be appropriate to estimate realistic data on the fiber/matrix interfaces in composites.
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Finally, the deviation of matrix cracks within the BN fiber coatings was observed only when a weak interface had been created in the coating. From this viewpoint, single layered BN coatings which showed a limited sensitivity to delamination, were not found as efficient as the pyrocarbon ones. Acknowledgements—This work has been supported by the french ministry of education and research and SEP through a grant given to F. R. and by C.N.R.S.. The authors are grateful to R. Kerans for his assistance in the push-in tests, to M. Lahaye for his help in the AES experiments, to C. RobinBrosse for his contribution to the preparation of 2D composites, to R. Naslain for valuable discussions and to J. Forget for the preparation of the manuscript. REFERENCES 1. Evans, A. G., J. Am. Ceram. Soc., 1990, 73, 187. 2. Kerans, R. J., Hay, R. S., Pagano, J. and Parthasarathy, T. A., Am. Ceram. Soc. Bull., 1989, 68, 429. 3. Naslain, R., Comp. Interf., 1993, 1(3), 253. 4. Droillard, C. and Lamon, J., J. Am. Ceram. Soc., 1996, 79(4), 849. 5. Droillard, C., Elaboration et caracterisation de composites a` matrice SiC et a` interphase se´quence´e C/SiC. Ph.D. thesis, Univ. Bordeaux (France), 1993, no. 913. 6. Guillaumat, L. and Lamon, J., Int. J. Fract., 1996, 82, 297. 7. Lissart, N. and Lamon, J., Acta Metall., 1997, 45(3), 1025. 8. Rebillat, F., Lamon, J., Naslain, R., Lara Curzio, E., Ferber, M. K. and Besmann, T., J. Am. Ceram. Soc., 1998, 81(4), 965. 9. Pasquier, S., Lamon, J. and Naslain, R., Composites Part A, 1998, 29, 1157. 10. Bertrand, S., Germain, F., Pailler, R. and Lamon, J., in Proceedings of the 12th International Conference on Composite Materials (ICCM12), Paris, 5–9 July 1999, in press. 11. Rugg, K. L., Tressler, R. E. and Lamon J., J. Eur. Ceram. Soc., in press. 12. Droillard, C., Lamon, J. and Bourrat, X., Mat. Res. Soc. Symp. Proc., 1994 MRS Fall Meeting, Materials Research Society, 1995, 365, 371.
13. Rebillat, F., Guette, A., Naslain, R. and Robin-Brosse, C., J. Eur. Ceram. Soc., 1997, 17, 1403. 14. Kerans, R. J. and Parthasarathy, T. A., J. Am. Ceram. Soc., 1991, 74(7), 1585. 15. Hsueh, C. H., J. Am. Ceram. Soc., 1993, 76(12), 3041. 16. Lara Curzio, E. and Ferber, M. K., J. Mater. Sci., 1994, 29, 6152. 17. Lara Curzio, E., Ferber, M. K. and Lowden, R. A., Ceram. Sci. Eng., 1994, 15(5), 989. 18. Kerans, K. J., Rebillat, F. and Lamon, J., J. Am. Ceram. Soc., 1997, 80(2), 506. 19. Rebillat, F., LCTS internal report, 1992. 20. Villeneuve, J. F., Mocaer, D., Pailler, R., Naslain, R. and Olry, P., J. Mater. Sci., 1993, 28, 1227. 21. Lamon, J., Lissart, N., Rechiniac, C., Roach, D. M. and Jouin, J. M., Ceram. Eng. Sci. Proc., 1993, Sept–Oct, 1115. 22. Lamon, J., Rechiniac, C., Lissart, N. and Corne, P., in Proceedings of the Fifth European Conference on Composite Materials (ECCM5), ed. A. R. Bunsell, J. F. Jamet and A. Massiah. EACM, Bordeaux, 1992, p. 895. 23. Corne, P., Rechiniac, C. and Lamon, J., in Proceedings JNC 8, ed. O. Allix, J. P. Favre and P. Ladeveze. AMAC, Paris, 1992, p. 213. 24. Lamon, J., Rebillat, F. and Evans, A. G., J. Am. Ceram. Soc., 1995, 78(2), 401–405. 25. Naslain, R., Ceram. Trans., 1995, 58, 23. 26. Prouhet, S., Cinetique de la CVD du nitrure de bore dans le systeme BF3–NH3–Ar. Application aux composites SiC/SiC a` interphase BN. Ph.D. thesis, University of Bordeaux I, 1991, no. 662. 27. Rebillat, F., Guette, A., Robin-Brosse, C. and Naslain, R., J. Am. Ceram. Soc., 1994, 77(3), 649. 28. Naslain, R., Dugne, O., Guette, A., Sevely, S., RobinBrosse, C., Rocher, J. P. and Cotteret, J., J. Am. Ceram. Soc., 1991, 74, 2482. 29. Rebillat, F., Lamon, J., Naslain, R., Lara Curzio, E., Ferber, M. K. and Besmann, T., J. Am. Ceram. Soc., 1998, 81(9), 2315. 30. Marshall, D. B., Cox, B. N. and Evans, A. G., Acta Metall., 1985, 33(11), 2013. 31. Aveston, J., Cooper, A. and Kelly, A., in Conference Proceedings of the National Physical Laboratory. IPC Science and Technology Press, Surrey, UK, 1971, p. 15.
THE CONCEPT OF A STRONG INTERFACE APPLIED TO SiC/SiC COMPOSITES WITH A BN INTERPHASE F. REBILLAT, J. LAMON† and A. GUETTE Laboratoire des Composites Thermostructuraux, UMR 5801, CNRS-SNECMA-CEA-UB1, 3, alle´e de la Boe´tie, 33600 Pessac, France
Abstract—Strong interfaces have been shown to allow improvement of the mechanical properties of ceramic matrix composites (CMC). The concept of a strong interface has been established in SiC/SiC composites with pyrocarbon (PyC) or multilayered (PyC/SiC) fiber coatings (also referred to as interphases). The present paper reports an attempt directed at applying the concept of a strong interface to SiC/SiC composites with a BN coating (referred to as SiC/BN/SiC). Fiber bonding and frictional sliding were investigated by means of push-out tests performed on 2D-composites as well as on microcomposite samples, and tensile tests performed on microcomposites. The stress–strain behavior of the SiC/BN/SiC composites and microcomposites is discussed with respect to interface characteristics and location of debonding either in the coating or in the fiber/coating interface. 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Interphase; Interface; Composites
1. INTRODUCTION
The fiber–matrix interfacial domain is a critical part of composites because load transfers from the matrix to the fiber and vice versa must occur through the interface. Therefore, it exerts a profound influence upon the mechanical behavior and the lifetime. Thus, it may be expected that composites could be tailored as a function of end use applications through optimization of interfaces. In fiber-reinforced ceramic composites, most authors promote the concept of weak interfaces to increase fracture toughness. The major contribution to toughness is attributed to crack bridging and fiber pull-out [1, 2]. Weak interfaces are detrimental to composite strength. A high strength requires efficient load transfers which are obtained with strong interfaces. This implies short debond cracks and/or significant sliding friction. These latter requirements, to be met for strong composites, are therefore incompatible with the former ones for tough composites, if toughening is based solely upon the above mentioned weak interface-based mechanisms. Fiber/matrix bonding results from diffusion or chemical reactions (chemical bonding) or from fiber clamping by residual stresses induced by thermal
† To whom all correspondence should be addressed. Tel.: ⫹33-556-844-703; fax: ⫹33-556-841-225. E-mail address: [email protected] (J. Lamon)
expansion mismatch. Fiber/matrix interfaces in the most advanced ceramic matrix composites consist of a thin coating layer (less than 1 µm thick) of one or several materials deposited on the fiber (interphase). Recently, SiC/SiC composites with strong interfaces have been developed. The coating/fiber bond was significantly stronger when fibers had been previously treated [3–5]. Features of the mechanical behavior of SiC/SiC composites with strong fiber/coating interfaces has been examined in several papers [4, 6–10]. Experiments as well as models have demonstrated that a strong interface is beneficial to the strength, the toughness, the lifetime and the creep resistance [4, 6– 11]. By contrast, weak interfaces are shown to be detrimental. The concept of strong interfaces has been established on SiC/C/SiC composites with PyC and multilayered (PyC/SiC) fiber coatings. In the present paper, it is applied to SiC/BN/SiC composites with boron nitride fiber coatings. BN is foreseen to be an alternative fiber coating to improve the oxidation resistance of ceramic matrix composites at high temperature.
2. FEATURES OF STRONG INTERFACES VS WEAK INTERFACES
In order to properly introduce the concept of a strong interface, we recall first the basic features of interface phenomena in CMCs subject to an essentially tensile load. These phenomena influence the
1359-6454/00/$20.00 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 0 ) 0 0 2 4 7 - 0
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mechanical response of composites described by the stress–strain curve. Fiber debonding results from the deflection of the cracks that initiate in the matrix (Fig. 1). Then sliding of the fiber debonded in the interface determines the load transfers from the fiber to the matrix and vice versa. The fiber sliding is influenced by the misfit strain, the associated radial component of the thermally induced residual stress-field, surface roughness and debond length. Weak interfaces debond easily. A single long debond crack is located at the surface of the fibers in those composites exhibiting weak interfaces (adhesive failure type, Fig. 1). As a consequence of small interface shear stresses, load transfers through the debonded interfaces are poor. The matrix becomes subjected to lower stresses and the volume of matrix that may experience further cracking is reduced by the presence of long debonds. The cracks are generally widely opened, whereas the crack spacing distance at saturation as well as the pull out length tend to be rather long (>100 µm). Toughening results essentially from sliding friction along the debonds. However, due to poor load transfers and long debonds, the fibers carry most of the load, which reduces the composite strength. The corresponding tensile stress–strain curve exhibits a short curved domain limited by a stress at matrix cracking saturation which is significantly smaller than ultimate strength (Fig. 2). In the presence of stronger fiber/coating bonds, the matrix cracks are deflected within the coating
Fig. 2. Typical tensile stress–strain behaviors measured on 2D SiC/SiC composites possessing PyC based interphases and fabricated from untreated or treated Nicalon (ceramic grade) fibers: (a) strong fiber/coating interfaces and (b) weak fiber/coating interfaces.
(cohesive failure type, Fig. 1), into short and branched multiple cracks [4, 12]. Short debonds as well as improved load transfers allow further cracking of the matrix via a scale effect [6, 7] leading to a higher density of matrix cracks (which are slightly opened). Sliding friction within the coating as well as multiple cracking of the matrix increase energy absorption, leading to toughening. Limited debonding and improved load transfers reduce the load carried by the fibers, leading to strengthening. The associated tensile stress strain curve exhibits a wide curved domain and the stress at matrix cracking saturation is close to ultimate failure (Fig. 2). Table 1 gives various values of the interfacial shear stresses measured using various methods on SiC/SiC composites with PyC-based fiber coating. It can be noticed that the interfacial shear stresses range between 10 and 20 MPa for the weak interfaces whereas they are larger than 100–300 MPa for the strong interfaces. Additional data can be found in [4, 7, 8, 24, 29]. 3. SiC/BN/SiC COMPOSITES: TESTING METHODOLOGY AND MICROSTRUCTURAL ANALYSES
3.1. Specimen preparation
Fig. 1. Schematic diagram showing crack deflection when the fiber coating/interface is (a) strong or (b) weak.
SiC/BN/SiC microcomposites and woven composites were manufactured via chemical vapor infiltration [13]. They were reinforced with either asreceived or treated (proprietary treatment, SNECMA/SEP, Bordeaux) SiC Nicalon fibers (NL 202 grade). The SiC/BN/SiC microcomposites consist of a single fiber (15 µm diameter), coated with a boron nitride layer (0.3–0.9 µm thick) and a SiC matrix deposited by CVD. They are representative of their counterparts in the 2D woven composites, since they are produced using identical chemical vapor deposition conditions [13]. A single or a bilayered BN fiber coating was deposited from a BF3, NH3, Ar gas mixture (Table
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4611
Table 1. Interfacial shear stress (MPa) measured using various methods on 2D-SiC/SiC composites with PyC based fiber coatings and reinforced with either as-received or treated fibers SiC/C/SiC SiC/(C/SiC)n/SiC
Untreated fibers 2D woven Microcomposites Minicomposites 2D woven
Treated fibers 2D woven
Interphase
Crack spacing [30]
Crack spacing [31]
Tensile tests (hysteresis loops) [24]
PyC (0.1) PyC(0.1) PyC(0.1) PyC(0.5) (PyC/SiC)2 (PyC/SiC)4
12
8
0.7 3 21–115 4 2 9
PyC(0.1) PyC(0.5) (PyC/SiC)2 (PyC/SiC)4
203
140
190 370 150 90
Tensile tests (curved domain) [7, 24]
Push-out tests (curved domain) [8, 29]
Push-out tests (plateau) [8, 29]
14–16 31 28
12–10 19.3 12.5
4–20 40–80
165–273 100–105 133 90
2). The selected processing conditions have been shown to cause minimum damage to the fibers and to improve adhesion of the BN coating onto the fibers, and the microstructure [13]. In the bi-layered coating (referred to as BN4, Table 2 and Fig. 3), the first sublayer on the fiber is BN2 type (poorly crystallized), whereas the second one is BN1 type (highly crystallized). Processing of BN2 involved the less aggressive gaseous phase, which led to a better contact between the fiber and the coating. The processing conditions of BN1 were found to be aggressive against the fibers [13]. 3.2. Push-out tests Fiber bonding and frictional sliding in the 2DSiC/BN/SiC composites were investigated by means of single fiber push-out tests [14–17]. 500 µm thick wedges were prepared using standard metallographic techniques. An interfacial test system (designed by ONERA, France) was used. The load was applied to the top of the fiber using a flat bottom diamond cone (at a constant displacement rate of 0.1 µm/s). The interface characteristics were extracted from the experimental stress–fiber end displacement curves by fitting the push-out model of Hsueh [15], as discussed in a previous paper [16]. Only a few push-out experiments could be carried out on the microcomposites owing to the difficulties involved in microcomposite handling, preparation and testing. Parallel-faced strips were cut out of the microcomposites which had been previously embedded in glass [18] (microcomposites 2) or in a ceramic cement [19] (microcomposites 2 and 4). A first series
Fig. 3. Images of BN coatings: (a) TEM-image and DEASpicture of the BN1 coating showing the three-dimensional ordered hexagonal structure; (b) SEM-image of bi-layered BN interphase (BN4) in a SiC/BN/SiC microcomposite showing poorly ordered BN2 layer and the BN1 layer with a threedimensional ordered hexagonal structure.
Table 2. Main characteristics of the BN fiber coatings [13] Batch
1 2 4
BN coating Number of Degree of BN layers crystallization BN1 BN2 BN4
1 1 2
High Low BN1⫹BN2
Coating adhesion Strong Weak Strong
of push out tests on microcomposites 2 used a Vickers diamond probe [18]. A correction for indentor displacement was done. Most of the tests were pushin rather than push-out tests. Push-out could not be achieved on the specimens with a thickness exceeding 200 µm. A number of push-in curves was unusable
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for analysis in that they exhibited features that were inconsistent with the model. A second series of pushin tests on microcomposites 2 was then performed on thicker samples (290 µm) using a flat-bottomed cone. Push-out tests were successfull on samples of microcomposites 4 (thickness ⬇190 µm). The interfaces characteristics were extracted by fitting the Hsueh’s model [15] to the push-in curves or to the curved domain and to the plateau of the push-out curves. 3.3. Tensile tests Five microcomposites per batch were tested in tension by using a specific table-model testing machine designed and developed for fiber testing [20]. The single fiber tensile test procedure based on window frames with appropriate gauge lengths (generally 10 mm) was employed [21, 22]. The microcomposites were loaded up to failure, either monotonically or with unloading–reloading cycles at a low strain rate (0.1% mn⫺1). The interface characteristics including the shear stress (t), the debond energy (Gic) and the debond length (ld) were extracted from the stress–strain curves [22–24] and from hysteresis loops on unloading–reloading [23, 24]. Independent models were used in order to assess the results. These models are referred to as LRLC, CLR and LRE according to authors’ names [22–24]. They derive from modelling the tensile load–displacement behavior (LRLC model) or the hysteretic behavior (CLR and LRE models) of microcomposites: the CLR model determines the energy dissipated in the friction phenomena whereas the LRE one determines the crack opening displacement during unloading–reloading cycles. After ultimate failure, the microcomposites were examined using scanning electron microscopy (SEM). The composition of the surface of fibers was determined from Auger electron spectroscopy (AES) depthprofile analyses of the pulled-out fibers. The tensile tests on the SiC/BN/SiC woven composites (three test specimens per batch) were performed at a constant strain rate of 0.05% min⫺1. Deformations were measured using an extensometer (gauge length 25 mm). The dimensions of the test specimens were as follows: thickness 3 mm, width 8 mm, length 100 mm. 4. RESULTS
4.1. Tensile tests on the SiC/BN/SiC microcomposites 4.1.1. Stress–strain curves. Most of the tensile stress–strain curves (Fig. 4) exhibited a curved domain over a wide range of deformations (0.2– 0.9%), and rather large strains-to-failure up to 1.2% (Table 3). However, most of the microcomposites with treated fibers essentially experienced premature failures. A plateau-like behavior was observed for
Fig. 4. Tensile stress–strain curves measured on the SiC/BN/SiC microcomposites reinforced with: (a) as-received fibers and (b) treated fibers.
microcomposites 1 and 4 reinforced with as-received fibers, suggesting the presence of rather weak fiber/matrix interactions, short debonds and small densities of matrix cracks at saturation when comparing with the microcomposites 2 which exhibit a widely curved stress–strain behavior up to ultimate failure and larger stresses. Saturation of matrix cracking generally occurred at rather large deformations (>0.6%). 4.1.2. SEM fractography of microcomposites. The numbers of matrix cracks identified on the microcomposites after ultimate failure were generally comparable with the numbers of load drops or of slope decreases on the force–displacement curves (Table 3). The higher density of matrix cracks was observed in microcomposites 2 (Table 3). In the microcomposites reinforced with untreated fibers debonding was observed mainly at the fiber/BN interface. The free surface of untreated fibers, on which the BN interphase was deposited, is at least partly made of silica [3]. The resulting fiber/BN interface has been reported to correspond to a very weak bond [3, 13, 25, 26]. Similar features have been observed on SiC/C/SiC composites with a fiber coating of anisotropic PyC [5]. In the microcomposites reinforced with treated fibers, debonding was detected in the BN coating only in microcomposites 4 with a bi-layered BN coating, which indicates that the weakest link is now located in the interface between the BN sublayers. As
REBILLAT et al.: SiC/SiC COMPOSITES
4613
Table 3. Main features of the stress–strain curves for the SiC/BN/SiC microcomposites and 2D woven composites
Interphase
Microcomposites
Composites
a b
BN1 BN2 BN4 BN1 BN2 BN4
Interphase thickness (µm)
0.28 0.27 0.29 0.5 0.3 0.5
Vf
0.42 0.76 0.47 0.40 0.40 0.40
Failure stress (MPa)
Failure strain (%)
Sa
Ta
S
792 1368 970 220 210 200
680 1813 670 32 110 210
0.55 0.85 0.99 0.58 0.38 0.5
T 0.18 1.27 0.2 0.06 0.11 0.064
Number of cracks at saturationb S
T
9 55 14
2 46 1
S⫽as-received fibers, T⫽treated fibers. Determined by SEM.
previously reported for Pyrocarbon fiber coatings [4, 5], treated fibers seem to give stronger fiber/BN bonds. However, the microcomposites with a single layer BN coating appear to be an exception to this rule, since the interface crack was detected at the fiber/BN interface. This was attributed to the presence of a weakly bonded sublayer of carbon that formed on the fibers [27]. 4.1.3. Auger electron spectroscopy analyses. AES depth-profile analyses of the pulled out fibers in microcomposites reinforced with untreated fibers, showed that the fiber surface is rich in free carbon. A layer enriched in carbon and oxygen (probably consisting of silica) is present under this carbon layer. Such a complex interfacial sequence has been already observed in 2D SiC/BN/SiC composites [26, 28]. The very thin carbon layer results from the attack of the fiber surface during BN processing [27]. 4.1.4. Extraction of interfacial properties from the stress–strain curve. The models provide comparable estimates of interfacial shear stresses for the microcomposites reinforced with untreated fibers (Fig. 5). A certain discrepancy may be observed for microcomposites 2. The interfacial shear stresses can be grouped into two distinct families (Fig. 5): t⬇5 MPa for microcomposites 4, tⱖ10 MPa for microcomposites 1 and 2, The debond energy estimates range between 1 and 8 J/m2 (Fig. 5). The interfacial characteristics determined for the microcomposites reinforced with treated fibers are shown on Fig. 6. The interfacial shear stresses obtained for microcomposites 1 and 2 are larger than those measured for the microcomposites reinforced with as-received fibers. A certain discrepancy is observed on the data extracted using the LRE model [24]: t ⫽ 400 MPa and Gic ⫽ 70 J/m2 seem to be overestimations although microcomposites 1 experienced a premature failure. The characteristics provided by the other models seem to be more realistic: 10⬍t⬍50 MPa, 0⬍Gic⬍7 J/m2 (Fig. 6).
Fig. 5. Interfacial characteristics estimated using various models for various BN interphases in SiC/BN/SiC microcomposites reinforced with untreated fibers.
4.2. Tensile tests on the 2D SiC/BN/SiC composites The stress–strain curves of the 2D SiC/BN/SiC composites also display a curved domain (Fig. 7). The strains-to-failure are smaller than those measured on the microcomposites (Table 3). They are close to 0.6% for composites 1 (reinforced with untreated fibers) and 4 (reinforced with as-received or treated fibers), whereas the other composites failed at deformations ⬍0.2%. 4.3. Push-out tests on the 2D SiC/BN/SiC composites 4.3.1. Composites reinforced with as-received fibers. The stresses to initiate and propagate the
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Fig. 6. Interfacial characteristics estimated using various models for various BN interphases in SiC/BN/SiC microcomposites reinforced with treated fibers.
debond crack as well as the interfacial shear stresses are much higher than those measured on 2D woven SiC/C/SiC composites with a PyC-based fiber coating and as received fibers [8, 29] (Table 4 and Fig. 8). As previously with the microcomposites, the largest t were obtained for composites 1 and 2. The debond stresses and the interfacial shear stresses thus appear to be insensitive to the conditions of BN processing. However, the effect of friction seems to be the most efficient in the composites 2, as indicated by the comparison of the respective values of the following specific parameters (Table 4): the applied maximum stress, the fiber-end displacement and the roughness amplitude. The smallest interfacial shear stress and roughness amplitude (A) were estimated for the bilayered BN coating (Table 4). SEM revealed the following interesting features: (i) in composites 1, the fiber surface was rather rough and the matrix surface showed a lot of pores, (ii) in composites 2, the slid surface was very smooth and debonding seemed to have occurred at the fiber/coating interface (as is usually observed in such composites [18, 26]), (iii) in composites 4, the debond crack was located in the interface between the BN sublayers (Fig. 9). 4.3.2. Composites reinforced with treated fibers. A large debond stress was estimated for composites 2 (sd ⫽ 2000 MPa, Table 4). It is much larger than
Fig. 7. Tensile stress–strain curves measured on the SiC/BN/SiC 2D composites reinforced with: (a) as-received fibers and (b) treated fibers.
those estimated for the composites reinforced with untreated fibers. Unfortunately, interfacial properties could not be extracted from the curved domain of the push-out curve, owing to bending of the sample as a result of the high load required to cause debonding. The interfacial shear stress that was extracted from the plateau is large (t⬇140 MPa). Smaller interfacial shear and debond stresses were obtained for composites 4 (Table 4). These interfacial shear stresses are close to those measured on the composites reinforced with untreated fibers. The magnitudes of debond stresses may be related to the respective stress–strain curves displayed by composites 2 and 4. The premature ultimate failure of composites 2 may correspond to the large sd value. The stress–strain behavior of composites 4, which is close to that observed for the composites reinforced with untreated fibers, may reflect the similarity in the respective interface characteristics t and sd. SEM revealed that debonding took place at the fiber/BN interface in composites 2. The contact between the coating and the fiber was not continuous, probably, as a result of the attack of the fiber by the gaseous phase. In composites 4, debonding occurred in the interface between the BN sublayers. Figure 9 shows that a BN sublayer remains bonded to the fiber. SEM examination of the fracture surface of 2D specimens tested in tension also showed a BN sublayer stuck on the fibers.
a
200 (10) 90 (15) 160 (50) 60 (15) 120 (10) 280 (10) 190 (10)
BN1 (Sa) BN2 (S) BN4 (S) BN2 (Ta) BN4 (T) BN2 (S) push-in BN4 (S) push-out
S⫽as-received fibers, T⫽treated fibers.
Microcomposites
Composites
Sample thickness (µm)
Interphase
1130 (200) 1250 (170) 1200 (330) 2000 (440) 740 (100) 1600 (400) 1100 (500)
sd (MPa) 3680 (1315) 3770 (530) 2720 (760) 2770 (350) 2670 (300) 5500 (900) 2500 (250)
smax (MPa)
0.9 (0.1) 1.6 (0.4) 0.6 (0.3)
1.35 (0.3) 1 (0.2) 1.1 (0.6)
Displacement (µm)
48 (9)
66 (9) 95 (20) 48 (9) 140 (26) 77 (9)
t (plateau) (MPa)
66 (5) 100 (40) 40 (15)
87 (22) 83 (15) 41 (13)
t (MPa)
0.06 (0.01) 0.11 (0.02) 0.04 (0.005)
0.11 (0.02) 0.09 (0.03) 0.09 (0.03)
m
68 (7) 50 (20) 45 (23)
32 (14) 45 (20) 16 (10)
A (nm)
120 (8) 110 (35) 110 (30)
120 (20) 95 (9) 105 (15)
ld (µm)
1.1 (0.4) 16 (8) 0.5 (0.4)
2.1 (1.1) 2.8 (1.4) 1.5 (1.9)
Gic (J/m2)
Table 4. Interfacial characteristics extracted from the push-out curves for the 2D-SiC/BN/SiC composites reinforced with as-received or treated fibers, and from push-in or push-out curves for the SiC/BN/SiC microcomposites reinforced with as-received fibers (standard deviation is given in brackets)
REBILLAT et al.: SiC/SiC COMPOSITES 4615
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REBILLAT et al.: SiC/SiC COMPOSITES
Fig. 8. Stress–fiber end displacement single fiber push-out curves measured on a SiC/BN/SiC microcomposite and a 2D SiC/BN/SiC composite (interphase BN4) reinforced with asreceived fibers.
The debond energy estimates fall within the range obtained from the tensile tests (Fig. 5). The debond stress in microcomposites 2 is quite high (sd ⫽ 1500 MPa), as well as the interfacial shear stress (t ⫽ 90 MPa), which indicates a high resistance to debonding. This may be related to the fact that the fiber could not be pushed out the matrix, even with thin samples. For microcomposites 4, a very good agreement can be noticed between the interfacial shear stresses extracted from respectively the curved domain and the plateau of the push out curves. As previously, the lower interfacial shear stress is obtained for this microcomposite with a double layer coating. Furthermore, push out of the fiber did not require a load as high as that (smax) applied during fiber pushing in the microcomposites 2. SEM showed that debonding occurred essentially at the fiber/coating interface. Generally, really smooth slid surfaces were observed. A certain roughness was detected. 5. DISCUSSION
Fig. 9. SEM micrographs showing protruding fibers after single fiber push-out tests performed on 2D-SiC/BN/SiC composites with an interphase made of two layers of boron nitride (BN4): (a) composite reinforced with as-received fibers and (b) composite reinforced with treated fibers.
4.4. Push-out and push-in tests on microcomposites The few push-out and push-in curves that were suitable for analysis exhibited the features usually observed with 2D composites (Fig. 8). The interface parameters extracted from these curves are summarized in Table 4. The interfacial shear stresses are larger than those extracted from the tensile stress–strain curves (Fig. 5). They are in excellent agreement with those estimated on the 2D-SiC/BN/SiC composites (Table 4).
The difficulty to perform push-out tests on thin samples becomes still worse with microcomposites. Large load transfers require very thin embedded lengths in order to permit push-out. The parallel-faced strip is not the most convenient test specimen because it is too fragile, but it remains the only usable sample geometry. During a push-in, unknown maximum debond length, damage of the fiber due to high compressive stresses, bending of the sample and diamond meeting with the matrix are factors that affect the stress–strain curves, and the validity of the extracted characteristics. Then, getting well polished thin strips with a thickness ⬍200 µm is rare. This explains why the push-out tests could be carried out only on microcomposites 4 which possessed the weakest interfaces. The less aggressive gaseous phase was used for processing the BN2 coating. This coating has been shown to adhere well to the fiber [13]. The resistance to propagation of the interface crack during the pushout tests appeared to be high, which suggests efficient load transfers under a tensile load. During the tensile tests, the composites 2 experienced a premature failure whereas the microcomposites 2 exhibited a high strain at saturation. Both features are compatible with efficient load transfers. Ultimate failure of 2D composites involves additionnal phenomena and it is affected by variability in fiber strength degradation during processing. The debond stresses and the interfacial shear stresses are larger than those obtained for composites and microcomposites 4 with a bi-layered coating, and they are comparable to those estimated for composites 1 with an interphase processed in more aggressive conditions. As previously mentioned, the propagation of interface cracks during the push-out tests was easier for composites 1 than
REBILLAT et al.: SiC/SiC COMPOSITES
for composites 2. This result is also indicated by the stress–strain tensile curves of microcomposites 1 which display a plateau-like feature when comparing with those of microcomposites 2. This suggests that the fiber BN1 bond is weaker. The plateau-like feature was not observed with the 2D composites. Furthermore, microcomposites 1 failed at a smaller deformation, which may be attributed to the degradation of the fiber during BN processing. The interface between the BN sublayers in composites 4 (with as-received or treated fibers), and in microcomposites 4 with treated fibers, was found to be comparatively weaker. The fact that the associated interfacial shear stresses are the smallest is not reflected by the stress–strain curves of the 2D composites nor of the microcomposites (premature failure). The interfacial shear stresses are not distinct enough to influence the stress–strain behavior. This conclusion agrees with predictions [6, 7]. With treated fibers, the fiber/coating bond was strengthened, as shown by the estimated interfacial shear stresses and the debond stresses. In the composites 2 reinforced with treated fibers, sd (which may be considered to be commensurate with the interface bond strength) and t were significantly increased. In the composites 4 with bi-layered fiber coatings, debonding occurred in the interface between the BN sublayers. Furthermore, the corresponding shear and debond stresses are small when comparing with those pertinent to the other SiC/BN/SiC composites for which debonding took place in the fiber/BN coating interface. It is particularly interesting to compare composites 2 and 4 for which a BN2 layer is on the fiber. It may be considered that the strength of the fiber/BN2 bond was the same in both composites. Therefore it may be concluded that the interface between the BN sublayers was weaker than the fiber/BN2 bond. The tensile behavior of the 2D SiC/BN/SiC composites may be considered to agree with the interface characteristics measured using push-out tests, since close interfacial shear stresses were obtained for the composites that exhibited a comparable stress–strain behavior, whereas the largest interfacial shear stresses were observed for the composites 2 with treated fibers that experienced a premature failure. This conclusion can be also drawn for the microcomposites reinforced with as-received fibers. The interfacial shear stresses are in a good correlation with the stress–strain behavior, except for microcomposites 1. Thus, the largest interfacial shear stresses were determined for microcomposites 2 which possessed the highest density of matrix cracks at saturation (Table 3) and gave the most pronounced curvature in the stress–strain curve. However, the thermally-induced residual stresses, resulting from the thermal expansion mismatch between the fibers and the matrix, exert a certain influence. They are increased by a low volume fraction of matrix, which may explain discrepancies in the stress–strain
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behavior. The microcomposites 2 possessed the thinnest matrix layer (Table 3). Comparable interfacial characteristics and an identical tensile stress–strain behavior were obtained for earlier SiC/BN/SiC composites with as-received Nicalon fibers but a different boron nitride [26]. A difference in ultimate failure can be noticed when comparing the microcomposites and the composites. This difference must be attributed to the mechanisms involved. As previously mentioned, the ultimate failure of 2D composites involves fiber interactions and the individual break of the weakest fibers prior to instability. The ultimate failure of microcomposites is thus dictated by the filament strength (which exhibits a wide statistical distribution), whereas that of 2D composites is determined by fiber tow strength (which exhibits a limited scatter). 6. CONCLUSION
In the SiC/BN/SiC microcomposites reinforced with untreated fibers, debonding occurred between the fiber free surface and the coating. The fiber/BN interface is the weakest bond in the interfacial sequence. The presence of a very thin layer of carbon between a SiO2–C mixture layer and the BN coating was detected by AES depth profile analyses. This sublayer seems to affect the interface bond. When the BN coating was deposited on treated fibers, the fiber/coating interface was stronger, but crack deflection did not occur within the BN coating. Only bi-layered BN coatings with a weak interface between the sublayers experienced crack deviation within the coating. The interface between the BN sublayers becomes the weakest link in the interfacial sequence between the fiber and the matrix. The micromechanics based models used for the analysis of the stress–strain behavior of the microcomposites provided comparable interfacial shear stresses, despite a certain scatter in some cases. A certain discrepancy was observed for the interface fracture energy estimates, which fell within a range of small values (1–10 J/m2). Although there might be some uncertainly in the t data estimated from tensile tests on microcomposites, they indicated stronger fiber/matrix bonds in the microcomposites fabricated with treated fibers. The data determined using the Hsueh’s pushout model were about one order of magnitude larger than those extracted from the tensile stress–strain behavior of microcomposites. The origin of this discrepancy has not been elucidated in the paper. The tensile stress/strain curves determined with microcomposites were similar to those obtained for the 2D composites. Tensile tests on microcomposites may be regarded as an interesting approach to interface design. Despite the difficulties to perform the push-out tests on microcomposites, this sample geometry seems to be appropriate to estimate realistic data on the fiber/matrix interfaces in composites.
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Finally, the deviation of matrix cracks within the BN fiber coatings was observed only when a weak interface had been created in the coating. From this viewpoint, single layered BN coatings which showed a limited sensitivity to delamination, were not found as efficient as the pyrocarbon ones. Acknowledgements—This work has been supported by the french ministry of education and research and SEP through a grant given to F. R. and by C.N.R.S.. The authors are grateful to R. Kerans for his assistance in the push-in tests, to M. Lahaye for his help in the AES experiments, to C. RobinBrosse for his contribution to the preparation of 2D composites, to R. Naslain for valuable discussions and to J. Forget for the preparation of the manuscript. REFERENCES 1. Evans, A. G., J. Am. Ceram. Soc., 1990, 73, 187. 2. Kerans, R. J., Hay, R. S., Pagano, J. and Parthasarathy, T. A., Am. Ceram. Soc. Bull., 1989, 68, 429. 3. Naslain, R., Comp. Interf., 1993, 1(3), 253. 4. Droillard, C. and Lamon, J., J. Am. Ceram. Soc., 1996, 79(4), 849. 5. Droillard, C., Elaboration et caracterisation de composites a` matrice SiC et a` interphase se´quence´e C/SiC. Ph.D. thesis, Univ. Bordeaux (France), 1993, no. 913. 6. Guillaumat, L. and Lamon, J., Int. J. Fract., 1996, 82, 297. 7. Lissart, N. and Lamon, J., Acta Metall., 1997, 45(3), 1025. 8. Rebillat, F., Lamon, J., Naslain, R., Lara Curzio, E., Ferber, M. K. and Besmann, T., J. Am. Ceram. Soc., 1998, 81(4), 965. 9. Pasquier, S., Lamon, J. and Naslain, R., Composites Part A, 1998, 29, 1157. 10. Bertrand, S., Germain, F., Pailler, R. and Lamon, J., in Proceedings of the 12th International Conference on Composite Materials (ICCM12), Paris, 5–9 July 1999, in press. 11. Rugg, K. L., Tressler, R. E. and Lamon J., J. Eur. Ceram. Soc., in press. 12. Droillard, C., Lamon, J. and Bourrat, X., Mat. Res. Soc. Symp. Proc., 1994 MRS Fall Meeting, Materials Research Society, 1995, 365, 371.
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