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Carbon-fiber-reinforced C-SiC binary matrix composites
Carbon, Vol. 33, No. 4, pp. 441-447, 1995 Copyright 0 1995 Elsevier Science Ltd
Pergamon
Printed in Great Britain. All rights reserved OOOS-6223/95 $9.50 + .OO
0008-6223(94)00169-3
CARBON-FIBER-REINFORCED C-Sic MATRIX COMPOSITES Institute
WENCHUAN of Metal Research,
BINARY
LIU, YONGLIANG WEI, and JINGYI DENG Academic Sinica, 72 Wenhua Road, 110015 Shenyang,
China
(Received 30 August 1994; accepted in revised form 10 November 1994) Abstract-C-Sic binary matrix composites, made from a porous C-C preform infiltrated with SIC, are mechanically characterized under tensile, compressive, flexural, and shear loading. The results of these tests show that the materials exhibit non-brittle fracture and have excellent strength. The oxidation behavior of the materials has been studied. Measured mass loss data from the C-Sic binary matrix composites indicated lower mass loss rate compared with Sic-coated C-C composites. The microstructures were investigated by X-ray diffraction, scanning electron microscope, and polarized optical microscope. The properties of C-SIC binary matrix composites are discussed in relation to the microstructures.
Key Words-Binary
matrix composites, oxidation, silicon carbide, CVI.
1. INTRODUCTION
come the problem of thermal mismatches between the C-C matrix and the coating layer. The aim of the present investigation is to discuss the binary matrix concept in composites processed according to the CVI technique, the emphasis being on C-Sic binary matrix composites. In this paper the mechanical and thermal properties of C-SIC binary matrix composites are also discussed in relation to the microstructure of matrices.
C-C composites are attracting considerable interest as materials for aerospace and high-temperature industrial application. This is due to their light weight and retention of high strength and stiffness at high temperatures[ 11. However, despite the excellent properties of C-C composites, a serious drawback is that carbon in any form will react with oxygen, burning away rapidly at temperatures as low as 5OO”C, so C-C composites are commonly used in short-time, high-temperature applications such as rocket propulsion components, re-entry thermal protectors, aircraft brakes, etc. Much effort is being expended to protect C-C composites against air-oxidation for a wide range of applications[2]. Many approaches are being considered to protect carbon from oxidation. These approaches fall into two categories[3]: (1) the addition of inhibitor to slow down the reaction rate, and (2) the use of oxygendiffusion barriers. At present, great interest is being focused on the possibility of coating C-C composites with protective layers that will be impertneable to oxygen at 1000°C and above. Unfortunately, because of thermal expansion difference and poor wetting properties, most ceramic materials do not form satisfactory coatings on carbon composites. Layers, for example, of Sic applied by CVD to C-C composites rapidly develop cracks in the SIC coating during cooling. Other CVD coating such as TiC and B,N,, when applied to C-C composites, also give incomplete oxidation protection[4]. In order to improve the crack resistance in coatings and their poor wetting properties, carbonfiber-reinforced C-SIC binary matrix composites were introduced by Naslain and co-workers in 1980[5]. The carbon matrix has been partially replaced by other refractory material, for example, a 2D C-C porous preform made from a stack of carbon fabrics consolidated by a small amount of pyrocarbon have been successfully densified according to the CVI technique by Sic. The concept of in-depth deposition of Sic rather than a simple overcoating was applied to over-
2. EXPERIMENTAL 2.1
CVZ technique
Chemical Vapor Infiltration (CVI) is an important technique for fabrication of carbon and ceramic matrix composites. There are two principal variations of CVI in current practice. In isothermal CVI, a gas mixture is introduced into a furnace containing one or more preforms. Reactant diffuses into the pore spaces of the preform, reacting on the fiber surface to form the matrix material. It suffers from limitation on component thickness, and requires very long processing time. This variation of the CVI process has gained the widest commercial application. The isothermal technique is illustrated in Fig. 1. In forced flow-thermal gradient CVI, the reactant gas mixture is forced to flow through the preform and an applied temperature gradient controls the progress of densification in order to avoid premature pore closure near the gas inlet. The forced flow-thermal gradient CVI is effective for thick components, and offers an order of magnitude improvement in processing time. Under development of ORNL for several years[6], this technique recently has seen commercial implementation. A schematic of the process is shown in Fig. 2. In the CVI processing of composites, the starting material is a porous fiber preform. In CVI, a solid is deposited within the pore network of a heated substrate as the result of a heterogeneous chemical reaction between gaseous species. In this work the C-Sic 441
W. Lm el al.
442
I m
The experimental set-up is illustrated in Fig. 3. The deposition apparatus used in this work was a SOKW graphite resistance furnace working zone dimension of 200 x 250 mm; see Fig. 1. The temperature of the furnace working zone was measured and controlled by a thermocouple. During the CVI process, the morphology and infiltration characteristics of C and SIC could be controlled by changing processing parameters, such as temperature, gas flow rate, gas concentration, deposition cycle, and so on. Table 1 shows the main parameters and their range of values considered in this experiment.
2.2 Specimens and testing procedures
I. Schematic
Fig.
of isothermal
CVI.
binary matrix composites were obtained according to a two-step CVI procedure. In the first step, the fibrous preform was consolidated by a small amount of pyrocarbon resulting from the pyrolysis of &Hz. At this stage, the partially densified material exhibits a large open porosity. In the second step, the porous C-C preform was progressively densified by SIC according to a CVI technique until the full densification of the preform. The infiltration of Sic was preformed from a gas mixture of an organosilane compound (e.g., CH3SiC13 or MTS), Ar, and HZ. It proceeds according to the overall equation: CH$Cl,
Fig. 2. Schematic
2
SIC + 3HCI.
of process utilizing gradient CVI.
forced
The preform used in this work consisted of a layup of layers of plain-weave carbon cloth maintained pressed together with a tooling or carbon felt. The composites were manufactured in panel form with average thickness of 6 mm and 8 mm for fabrics and felt, respectively. In this paper, these materials are identified as CVD/cloth-SiC and CVD/felt-Sic. All the samples used for tensile, compressive, flexUral, shear, thermal conductivity, thermal expansion, oxidation resistance, and specific heat measurements were cut from the homogenous blocks of the corresponding material and machined to size. Mechanical tests were conducted on a WDJ-10 universal testing machine. Since C-Sic binary matrix composites tend to behave in a rather brittle manner, the tests were run under displacement control. Tensile strength, strain, and modulus were measured in plane direction. Specimen size was 8 mm wide x 6 mm thick x 30 mm gauge length. The test was run at cross-head speed of 0.25-l mm/min. Compressive strength, strain, and modulus were measured on specimens 10 x 10 x 25 mm at a crosshead speed of 0.1-0.5 mm/min. Flexural strength and modulus were measured in three-point flexural on specimens 100 x 15 x 4 mm using a span: depth of 20: 1 at a cross-head separation of 1 mm/min. lnterlaminar shear strength was determined using specimens 30 x 6 x 4 mm at a span:depth ratio 5: 1 and a cross-head separation of 1 mm/min. In the interlaminar shear test, which employed a three-point flexural loading configuration, it was assumed that the
flow-thermal Fig. 3. Schematic
diagram
of CVI system
Carbon-fiber-reinforced
SIC composites
443
Table 1. Deposition parameters of carbon and Sic Rate of source gas (M3/h)
Method Isothermal Forced flowthermal gradienta
Mater.
Temp. (“C)
Pressure (Pa)
C2H, CH,SiCI, + Arb
C Sic
loo0 900-1000
300-700 300-loo0
0.02
Sic
900-1000 Exhaust gas: 700-2000 Inlet gas: 2ooo-6oooO
‘Isothermal CVI is used to deposit pyrocarbon aration of the C-SIC binary matrix composites. bAr gas saturated with CH,SiCl, vapor.
H,
0.04
0.01
0.06
and then Sic is infiltrated
3P
where p = load applied at the center of the beam, b = width of the beam, and h = depth of the beam. Linear thermal expansion tests have been conducted in an inert atmosphere, at 1300°C and 16OO”C, on samples 7 mm diameter x 50 mm long, with carbon cloth layers parallel to the cylinder sample axis. Thermal diffusivity measurements were performed on specimens 10 mm diameter x 2 mm thickness, according to the flash method with, carbon cloth layers perpendicular to the cylinder axis and heat flux direction. In this technique, one face of the sample was illuminated during a very short time with a laser beam; the thermal diffusivity was calculated from the recording of the change of the temperature of the opposite face vs time. The thermal conductivity is given by the following equation: CY&,,
where K = thermal conductivity, o = thermal diffusivity, p = density of material, and C, = specific heat. The isothermal oxidation behavior of the C-Sic binary matrix composites at an elevated temperature was carried out in a muffle. Specimen size were 50 x 50 x 6 mm. The temperature was raised to the testing point, and specimen was loaded directly to the hot zone. The weight change was measured after it was taken out from the furnace and cooled down to room temperature. In the present work, the specimens were oxidized at 540°C and 1260°C for 10 h in air environment, respectively. Generally, oxidation rates were calculated based on the weight loss of specimen, the total oxidation time, and the initial surface area of the specimen:
Deposition cycle Reinforcement
0.05
4bh
K =
Ar
0.004-0.01
shear stress r was parabolically distributed through the depth of the specimen with the maximum occurring at the center of the beam given by
7=-,
Rate of dilute gas (M3/h)
Carbon cloth lay-up or felt
by forced
M=---
flow-thermal
(h)
Density
30-50 700
0.9-I .o 1.6
130
1.8
gradient
w-
w-2
ST
’
during
prep-
where M = mass loss rate, W, = original weight of the specimen, W, = weight after oxidation specimen, S = surface area of the specimen, and T = exposure time interval. In order to simulate a more realistic application, a thermal cycling test was conducted. The specimen, nominally 100 x 100 x 6 mm, was plunged into the furnace directly from room temperature. The increase in temperature from room temperature to 1250°C was completed in one second, then held at 1250°C for 5 min, and finally cooling to room temperature. This processing cycle was repeated 30 times. In addition, the candle quartz insulation test was performed on specimens 100 x 100 x 6 mm. The specimen was exposed to a heat flux of 120 kcal/msec, and after 5 l/2 min the temperature of the backside was measured. The morphology of C-Sic binary matrix composites was observed by using the polarized optical microscope and scanning electron microscope. Microstructure analyses were conducted by X-ray.
3. RESULTS AND DISCUSSION Chemical Vapor Infiltration has been utilized by previous researchers to form the matrix of fiber-reinforced ceramic composites[7-101. The process offers near-net-shape fabrication of structural components using materials that cannot be processed using the methods developed for polymers and metals, and has been implemented in commercial-scale production facilities in Europe and the United States. CVI is probably the most common method currently used for producing ceramic matrix composites, but without exception, processing times required for CVD infiltration have been extremely long. CVI processes induce highly controlled reactions in vapor-phase reactant within hot-wall hermetic chambers. Temperature, reactant partial pressure, and total pressure are the main
Llu
et al
(4 Fig. 4. The optical
micrograph
of C-Sic
binary
process variables with a specific set of reactant. The foremost issue is one of finding conditions by which uniform depositions can be achieved at acceptably high rates throughout the entire volume of the preform. Conventional isothermal CVI processes are slow, and even thin parts require several weeks to several months to full infiltration. Even under these process conditions, severe density gradients still occur in composites for thickness greater than 3 mm. The forced CVI method under development at ORNL has resulted in improvements in processing times from weeks to several hours. As seen from the data in Table 1, the forced CVI method not only reduces the processing times, but also obtains higher infiltrated densities. This process is well suited for the fabrication of thickwalled composites of relatively simple shapes. Although isothermal CVI processes are slow, economical processes can be achieved through the use of large furnaces that contain multiple preforms, and this process is suited to the fabrication of thin-walled composites of complex shapes. The pyrocarbon and Sic matrices studied in this work were analysed by optical microscope using polarized light. Figure 4 shows a typical polarized optical micrograph of C-Sic binary matrix composites. As seen from Fig. 4b, there are two different sheaths moving outward from the carbon fiber core. The inner pyrocarbon layer has a rough laminar texture and the outer SIC layer has an isotropic texture. Systematic observations of the C-Sic interface in CVD/felt-SiC samples did not reveal the existence of microfissures. It seems that the cohesion between the fiber and carbon layer, as well as between the carbon and Sic composing the matrix was rather strong. However, in analyzing the specimens after they had failed, it was noticed that the weakest interface in these materials was the fiber-carbon interface, rather than Sic-C interface, as shown in Fig. 5. Table 2 presents mechanical and thermal property data of the C-Sic binary matrix composites. In prac-
matrix
composites
under
polarized
light
tical experiments, an important spread of the mechanical property data was observed partly due to a microstructure inhomogeneity in the initial C-C preforms and to the difference in composite density. Tests in tension at room temperature show a typical stressstrain curve with three regions, as shown in Fig. 6. There are: (a) linear stress-strain behavior before matrix cracking, (b) nonlinear region with increasing stress and strain where multiple matrix cracking occurred with possible crack deflection at the fibermatrix interface, and (c) a region of decreasing stress with increasing strain where fiber pullout occurred. In the absence of carbon matrix (i.e., the fibers directly in contact with SIC matrix), the material (C-Sic) exhibited low strength, low toughness, and brittle fracture behavior because the adhesion between the carbon fiber and SIC matrix was complete, so that no crack energy could be absorbed by delamination. These composites fractured by the propagation of a single, near-planar crack that passed through the matrix and fiber with no evidence of interface debonding, as
Fig. 5. SEM micrograph of CVD/cloth-SiC showing weak bond between carbon fiber and pyrocarbon layer.
Carbon-fiber-reinforced Table 2. Mechanical
and thermal
Properties
Unit
Density Tensile strength Tensile modulus Tensile failure strain Compressive strength
g/cm’ MPa GPa % MPa
Compressive
GPa
modulus
Compressive failure Flexura strength
Sic composites properties
Flexural modulus Interlaminar shear strength Mass loss rate 54O”C, 10 h 126O”C, 10 h Thermal expansion coeff. Thermal conductivity Candle quartz insulation (heat flux: 120 Kcal/m.s, temp. of back: 917-933”C, time: 5; min) Alternating temperature (5-min cycles at 1230-1250°C, interval of 5 mitt, repeated 30 times) _
of C-C-SIC
Direction
GPa MPa g/m2 .s g/m2.s x 1o-6 W/m.K
composites
CVD/cloth-SiC
CVD/felt-SiC
1.5-1.6 192-216 32.7-49.8 0.58-0.74 95-145 200-239 48-92 1.53-2.09 0.15-0.26 50.6-74.9 139-266 20.7-41.9 10.5-12.6
1.6 19-31
0.00151 0.00171 1.17-l .23 5.03-6.39 No change in appearance
0.05
II II II II _L II I II II I II II
% MPa
strain
445
Colour of surface becomes light grey
shown in Fig. 7. Consequently, in ceramic composites, the nature of the fiber-matrix interface is a key factor in determining strength and toughness. C-Sic binary matrix composites have a potentially better resistance to oxidation than coated composites. It was shown by our own research studies[ 1 l] that the mass loss rate of the C-SIC binary matrix composites was 0.00171 g/m3 s at 1260°C after 10 h and that of
Sic-coated C-C composites was 0.00351 g/m3 s under similar conditions. Most of the refractory materials used as protective layers were not satisfactory due to the microcracks induced by the thermal stress in the coating. Any coating used to protect the composites from oxidation must prevent the inward diffusion of oxygen and the outward diffusion of carbon. The coating must have good adherence to the substrate and
r
\ In
.\ \ \ I
0
0.2
0.4
0.6
0.8
1.0
1.2
1.4
1.6
1.8
2.0 (3%)
Fig. 6. Stress-strain
curve of CVD/cloth-Sic.
2.2
W. Liu et al.
Fig. 7. SEM micrograph of C-Sic showing between pyrocarbon and Sic.
strong
bond
little penetration. Because deposition of SIC occurs at high temperature, cooling the composite inevitably leads to microcracking the coating. Oxidation results from oxygen transport along the microcracks to the underlying C-C; in consequence, all oxidizable materials have been removed from within the sample and a hollow Sic shell has been left. The C-Sic binary matrix composites used in the present study was Sic matrix reinforced with carbon fibers separated by a carbon layer (i.e., the carbon fibers were first covered with a thin layer of pyrocarbon and then further densified by Sic). Consequently, the protective layer around each fiber consisted of an underlayer portion of pyrocarbon and overlayer portion of SIC. Owing to the small volume expansion of the carbon fiber the
Fig. 8. Typical
X-ray
diffraction
pattern
stress in the coating is reduced; hence, the cracks in the overlayer portion of Sic cannot occur. If the densified composite is machined to a desired configuration, the fiber ends will be exposed by machining. Oxidation appeared to lead to the propagation of longitudinal channels along the fiber axes and formed annular cavities. In the present study, a layer of Sic was applied to the exposed surface of the machined composites to provide a surface oxidation barrier. As a result, serious oxidation damage was not observed. From thermal cycling tests, it can be seen that C-Sic binary matrix composites show sufficient thermal shock resistance. The structural integrity of material is kept without splitting and spalling[l2]. X-ray diffractometry of samples throughout the thermal exposure testing revealed the SiOz phase present in the coating, as shown in Fig. 8. Under oxidizing conditions, at high temperature, the silicon carbide is converted to silicon dioxide, leading to the formation of a protective film of silica.
4. CONCLUSIONS
Chemical Vapor Infiltration has been successful in fabricating advanced ceramic composites. Carbonfiber-reinforced C-SIC binary matrix composites are of considerable interest because of their unique properties, such as excellent strength, sufficient thermoshock resistance, and good resistance to oxidation. A comparison of the data for different materials oxidation tests show that the mass loss rate of C-Sic binary matrix composites is lower than that of Sic-coated C-C composites. It is also noted that extensive fiber
of C-SIC binary
matrix
composites
after oxidation.
Carbon-fiber-reinforced pullout occurred in C-Sic binary matrix composites. The pyrocarbon interlayer between the carbon fiber and Sic matrix is believed to form a bond strong enough for load transfer, yet weak enough to debond readily and allow fiber pullout during crack propagation, leading to increases in strength and toughness.
REFERENCES
1. J. D. Buckely, Am. Ceram. Sot. Bull. 67, 364 (1988). 2. J. R. Strife and J. E. Sheehan, Am. Ceram. Sot. Bull. 67, 369 (1988). 3. K. L. Luthra, Carbon 26, 217 (1988). 4. D. W. McKee, Carbon 25, 551 (1987). 5. R. Naslain, P. Hagen Muller, F. Christin, L. Heraud,
Sic composites
447
and J. J. Choury, In Proc. 3rd Int. Conf. Composite Paris (1980). 6. D. P. Stinton, A. J. Caputo, and R. A. Lowden, Am. Materials,
Ceram. Sot. Bull. 65, 347 (1986). I. E. Fitzer and R. Gadow, Am. Ceram. Sot. BUD. 65, 326 (1986). 8. D. P. Stinton, T. M. Besmann, and R. A. Lowden, Am. Ceram. Sot. BUN. 67, 350 (1988). 9. Y. M. Chiang, J. S. Haggerty, R. P. Messmer, and C. Demetry, Am. Ceram. Sot. Bull. 68, 420 (1989). 10. L. J. Schioler and J. J. Stiglich, Jr., Am. Ceram. Sot. BUN. 65, 289 (1986). 11. W. Liu, S. Sun, and M. Li, Proceedings Carbon ‘90,
Paris (1990), p. 492. 12. W. Liu, S. Sun, M. Li, and Y. Wei, Proceedings Carbon ‘92, Essen, Germany, Paper G38, Deutsche Keramische Gesellschaft (1992), p. 741.
Pergamon
Printed in Great Britain. All rights reserved OOOS-6223/95 $9.50 + .OO
0008-6223(94)00169-3
CARBON-FIBER-REINFORCED C-Sic MATRIX COMPOSITES Institute
WENCHUAN of Metal Research,
BINARY
LIU, YONGLIANG WEI, and JINGYI DENG Academic Sinica, 72 Wenhua Road, 110015 Shenyang,
China
(Received 30 August 1994; accepted in revised form 10 November 1994) Abstract-C-Sic binary matrix composites, made from a porous C-C preform infiltrated with SIC, are mechanically characterized under tensile, compressive, flexural, and shear loading. The results of these tests show that the materials exhibit non-brittle fracture and have excellent strength. The oxidation behavior of the materials has been studied. Measured mass loss data from the C-Sic binary matrix composites indicated lower mass loss rate compared with Sic-coated C-C composites. The microstructures were investigated by X-ray diffraction, scanning electron microscope, and polarized optical microscope. The properties of C-SIC binary matrix composites are discussed in relation to the microstructures.
Key Words-Binary
matrix composites, oxidation, silicon carbide, CVI.
1. INTRODUCTION
come the problem of thermal mismatches between the C-C matrix and the coating layer. The aim of the present investigation is to discuss the binary matrix concept in composites processed according to the CVI technique, the emphasis being on C-Sic binary matrix composites. In this paper the mechanical and thermal properties of C-SIC binary matrix composites are also discussed in relation to the microstructure of matrices.
C-C composites are attracting considerable interest as materials for aerospace and high-temperature industrial application. This is due to their light weight and retention of high strength and stiffness at high temperatures[ 11. However, despite the excellent properties of C-C composites, a serious drawback is that carbon in any form will react with oxygen, burning away rapidly at temperatures as low as 5OO”C, so C-C composites are commonly used in short-time, high-temperature applications such as rocket propulsion components, re-entry thermal protectors, aircraft brakes, etc. Much effort is being expended to protect C-C composites against air-oxidation for a wide range of applications[2]. Many approaches are being considered to protect carbon from oxidation. These approaches fall into two categories[3]: (1) the addition of inhibitor to slow down the reaction rate, and (2) the use of oxygendiffusion barriers. At present, great interest is being focused on the possibility of coating C-C composites with protective layers that will be impertneable to oxygen at 1000°C and above. Unfortunately, because of thermal expansion difference and poor wetting properties, most ceramic materials do not form satisfactory coatings on carbon composites. Layers, for example, of Sic applied by CVD to C-C composites rapidly develop cracks in the SIC coating during cooling. Other CVD coating such as TiC and B,N,, when applied to C-C composites, also give incomplete oxidation protection[4]. In order to improve the crack resistance in coatings and their poor wetting properties, carbonfiber-reinforced C-SIC binary matrix composites were introduced by Naslain and co-workers in 1980[5]. The carbon matrix has been partially replaced by other refractory material, for example, a 2D C-C porous preform made from a stack of carbon fabrics consolidated by a small amount of pyrocarbon have been successfully densified according to the CVI technique by Sic. The concept of in-depth deposition of Sic rather than a simple overcoating was applied to over-
2. EXPERIMENTAL 2.1
CVZ technique
Chemical Vapor Infiltration (CVI) is an important technique for fabrication of carbon and ceramic matrix composites. There are two principal variations of CVI in current practice. In isothermal CVI, a gas mixture is introduced into a furnace containing one or more preforms. Reactant diffuses into the pore spaces of the preform, reacting on the fiber surface to form the matrix material. It suffers from limitation on component thickness, and requires very long processing time. This variation of the CVI process has gained the widest commercial application. The isothermal technique is illustrated in Fig. 1. In forced flow-thermal gradient CVI, the reactant gas mixture is forced to flow through the preform and an applied temperature gradient controls the progress of densification in order to avoid premature pore closure near the gas inlet. The forced flow-thermal gradient CVI is effective for thick components, and offers an order of magnitude improvement in processing time. Under development of ORNL for several years[6], this technique recently has seen commercial implementation. A schematic of the process is shown in Fig. 2. In the CVI processing of composites, the starting material is a porous fiber preform. In CVI, a solid is deposited within the pore network of a heated substrate as the result of a heterogeneous chemical reaction between gaseous species. In this work the C-Sic 441
W. Lm el al.
442
I m
The experimental set-up is illustrated in Fig. 3. The deposition apparatus used in this work was a SOKW graphite resistance furnace working zone dimension of 200 x 250 mm; see Fig. 1. The temperature of the furnace working zone was measured and controlled by a thermocouple. During the CVI process, the morphology and infiltration characteristics of C and SIC could be controlled by changing processing parameters, such as temperature, gas flow rate, gas concentration, deposition cycle, and so on. Table 1 shows the main parameters and their range of values considered in this experiment.
2.2 Specimens and testing procedures
I. Schematic
Fig.
of isothermal
CVI.
binary matrix composites were obtained according to a two-step CVI procedure. In the first step, the fibrous preform was consolidated by a small amount of pyrocarbon resulting from the pyrolysis of &Hz. At this stage, the partially densified material exhibits a large open porosity. In the second step, the porous C-C preform was progressively densified by SIC according to a CVI technique until the full densification of the preform. The infiltration of Sic was preformed from a gas mixture of an organosilane compound (e.g., CH3SiC13 or MTS), Ar, and HZ. It proceeds according to the overall equation: CH$Cl,
Fig. 2. Schematic
2
SIC + 3HCI.
of process utilizing gradient CVI.
forced
The preform used in this work consisted of a layup of layers of plain-weave carbon cloth maintained pressed together with a tooling or carbon felt. The composites were manufactured in panel form with average thickness of 6 mm and 8 mm for fabrics and felt, respectively. In this paper, these materials are identified as CVD/cloth-SiC and CVD/felt-Sic. All the samples used for tensile, compressive, flexUral, shear, thermal conductivity, thermal expansion, oxidation resistance, and specific heat measurements were cut from the homogenous blocks of the corresponding material and machined to size. Mechanical tests were conducted on a WDJ-10 universal testing machine. Since C-Sic binary matrix composites tend to behave in a rather brittle manner, the tests were run under displacement control. Tensile strength, strain, and modulus were measured in plane direction. Specimen size was 8 mm wide x 6 mm thick x 30 mm gauge length. The test was run at cross-head speed of 0.25-l mm/min. Compressive strength, strain, and modulus were measured on specimens 10 x 10 x 25 mm at a crosshead speed of 0.1-0.5 mm/min. Flexural strength and modulus were measured in three-point flexural on specimens 100 x 15 x 4 mm using a span: depth of 20: 1 at a cross-head separation of 1 mm/min. lnterlaminar shear strength was determined using specimens 30 x 6 x 4 mm at a span:depth ratio 5: 1 and a cross-head separation of 1 mm/min. In the interlaminar shear test, which employed a three-point flexural loading configuration, it was assumed that the
flow-thermal Fig. 3. Schematic
diagram
of CVI system
Carbon-fiber-reinforced
SIC composites
443
Table 1. Deposition parameters of carbon and Sic Rate of source gas (M3/h)
Method Isothermal Forced flowthermal gradienta
Mater.
Temp. (“C)
Pressure (Pa)
C2H, CH,SiCI, + Arb
C Sic
loo0 900-1000
300-700 300-loo0
0.02
Sic
900-1000 Exhaust gas: 700-2000 Inlet gas: 2ooo-6oooO
‘Isothermal CVI is used to deposit pyrocarbon aration of the C-SIC binary matrix composites. bAr gas saturated with CH,SiCl, vapor.
H,
0.04
0.01
0.06
and then Sic is infiltrated
3P
where p = load applied at the center of the beam, b = width of the beam, and h = depth of the beam. Linear thermal expansion tests have been conducted in an inert atmosphere, at 1300°C and 16OO”C, on samples 7 mm diameter x 50 mm long, with carbon cloth layers parallel to the cylinder sample axis. Thermal diffusivity measurements were performed on specimens 10 mm diameter x 2 mm thickness, according to the flash method with, carbon cloth layers perpendicular to the cylinder axis and heat flux direction. In this technique, one face of the sample was illuminated during a very short time with a laser beam; the thermal diffusivity was calculated from the recording of the change of the temperature of the opposite face vs time. The thermal conductivity is given by the following equation: CY&,,
where K = thermal conductivity, o = thermal diffusivity, p = density of material, and C, = specific heat. The isothermal oxidation behavior of the C-Sic binary matrix composites at an elevated temperature was carried out in a muffle. Specimen size were 50 x 50 x 6 mm. The temperature was raised to the testing point, and specimen was loaded directly to the hot zone. The weight change was measured after it was taken out from the furnace and cooled down to room temperature. In the present work, the specimens were oxidized at 540°C and 1260°C for 10 h in air environment, respectively. Generally, oxidation rates were calculated based on the weight loss of specimen, the total oxidation time, and the initial surface area of the specimen:
Deposition cycle Reinforcement
0.05
4bh
K =
Ar
0.004-0.01
shear stress r was parabolically distributed through the depth of the specimen with the maximum occurring at the center of the beam given by
7=-,
Rate of dilute gas (M3/h)
Carbon cloth lay-up or felt
by forced
M=---
flow-thermal
(h)
Density
30-50 700
0.9-I .o 1.6
130
1.8
gradient
w-
w-2
ST
’
during
prep-
where M = mass loss rate, W, = original weight of the specimen, W, = weight after oxidation specimen, S = surface area of the specimen, and T = exposure time interval. In order to simulate a more realistic application, a thermal cycling test was conducted. The specimen, nominally 100 x 100 x 6 mm, was plunged into the furnace directly from room temperature. The increase in temperature from room temperature to 1250°C was completed in one second, then held at 1250°C for 5 min, and finally cooling to room temperature. This processing cycle was repeated 30 times. In addition, the candle quartz insulation test was performed on specimens 100 x 100 x 6 mm. The specimen was exposed to a heat flux of 120 kcal/msec, and after 5 l/2 min the temperature of the backside was measured. The morphology of C-Sic binary matrix composites was observed by using the polarized optical microscope and scanning electron microscope. Microstructure analyses were conducted by X-ray.
3. RESULTS AND DISCUSSION Chemical Vapor Infiltration has been utilized by previous researchers to form the matrix of fiber-reinforced ceramic composites[7-101. The process offers near-net-shape fabrication of structural components using materials that cannot be processed using the methods developed for polymers and metals, and has been implemented in commercial-scale production facilities in Europe and the United States. CVI is probably the most common method currently used for producing ceramic matrix composites, but without exception, processing times required for CVD infiltration have been extremely long. CVI processes induce highly controlled reactions in vapor-phase reactant within hot-wall hermetic chambers. Temperature, reactant partial pressure, and total pressure are the main
Llu
et al
(4 Fig. 4. The optical
micrograph
of C-Sic
binary
process variables with a specific set of reactant. The foremost issue is one of finding conditions by which uniform depositions can be achieved at acceptably high rates throughout the entire volume of the preform. Conventional isothermal CVI processes are slow, and even thin parts require several weeks to several months to full infiltration. Even under these process conditions, severe density gradients still occur in composites for thickness greater than 3 mm. The forced CVI method under development at ORNL has resulted in improvements in processing times from weeks to several hours. As seen from the data in Table 1, the forced CVI method not only reduces the processing times, but also obtains higher infiltrated densities. This process is well suited for the fabrication of thickwalled composites of relatively simple shapes. Although isothermal CVI processes are slow, economical processes can be achieved through the use of large furnaces that contain multiple preforms, and this process is suited to the fabrication of thin-walled composites of complex shapes. The pyrocarbon and Sic matrices studied in this work were analysed by optical microscope using polarized light. Figure 4 shows a typical polarized optical micrograph of C-Sic binary matrix composites. As seen from Fig. 4b, there are two different sheaths moving outward from the carbon fiber core. The inner pyrocarbon layer has a rough laminar texture and the outer SIC layer has an isotropic texture. Systematic observations of the C-Sic interface in CVD/felt-SiC samples did not reveal the existence of microfissures. It seems that the cohesion between the fiber and carbon layer, as well as between the carbon and Sic composing the matrix was rather strong. However, in analyzing the specimens after they had failed, it was noticed that the weakest interface in these materials was the fiber-carbon interface, rather than Sic-C interface, as shown in Fig. 5. Table 2 presents mechanical and thermal property data of the C-Sic binary matrix composites. In prac-
matrix
composites
under
polarized
light
tical experiments, an important spread of the mechanical property data was observed partly due to a microstructure inhomogeneity in the initial C-C preforms and to the difference in composite density. Tests in tension at room temperature show a typical stressstrain curve with three regions, as shown in Fig. 6. There are: (a) linear stress-strain behavior before matrix cracking, (b) nonlinear region with increasing stress and strain where multiple matrix cracking occurred with possible crack deflection at the fibermatrix interface, and (c) a region of decreasing stress with increasing strain where fiber pullout occurred. In the absence of carbon matrix (i.e., the fibers directly in contact with SIC matrix), the material (C-Sic) exhibited low strength, low toughness, and brittle fracture behavior because the adhesion between the carbon fiber and SIC matrix was complete, so that no crack energy could be absorbed by delamination. These composites fractured by the propagation of a single, near-planar crack that passed through the matrix and fiber with no evidence of interface debonding, as
Fig. 5. SEM micrograph of CVD/cloth-SiC showing weak bond between carbon fiber and pyrocarbon layer.
Carbon-fiber-reinforced Table 2. Mechanical
and thermal
Properties
Unit
Density Tensile strength Tensile modulus Tensile failure strain Compressive strength
g/cm’ MPa GPa % MPa
Compressive
GPa
modulus
Compressive failure Flexura strength
Sic composites properties
Flexural modulus Interlaminar shear strength Mass loss rate 54O”C, 10 h 126O”C, 10 h Thermal expansion coeff. Thermal conductivity Candle quartz insulation (heat flux: 120 Kcal/m.s, temp. of back: 917-933”C, time: 5; min) Alternating temperature (5-min cycles at 1230-1250°C, interval of 5 mitt, repeated 30 times) _
of C-C-SIC
Direction
GPa MPa g/m2 .s g/m2.s x 1o-6 W/m.K
composites
CVD/cloth-SiC
CVD/felt-SiC
1.5-1.6 192-216 32.7-49.8 0.58-0.74 95-145 200-239 48-92 1.53-2.09 0.15-0.26 50.6-74.9 139-266 20.7-41.9 10.5-12.6
1.6 19-31
0.00151 0.00171 1.17-l .23 5.03-6.39 No change in appearance
0.05
II II II II _L II I II II I II II
% MPa
strain
445
Colour of surface becomes light grey
shown in Fig. 7. Consequently, in ceramic composites, the nature of the fiber-matrix interface is a key factor in determining strength and toughness. C-Sic binary matrix composites have a potentially better resistance to oxidation than coated composites. It was shown by our own research studies[ 1 l] that the mass loss rate of the C-SIC binary matrix composites was 0.00171 g/m3 s at 1260°C after 10 h and that of
Sic-coated C-C composites was 0.00351 g/m3 s under similar conditions. Most of the refractory materials used as protective layers were not satisfactory due to the microcracks induced by the thermal stress in the coating. Any coating used to protect the composites from oxidation must prevent the inward diffusion of oxygen and the outward diffusion of carbon. The coating must have good adherence to the substrate and
r
\ In
.\ \ \ I
0
0.2
0.4
0.6
0.8
1.0
1.2
1.4
1.6
1.8
2.0 (3%)
Fig. 6. Stress-strain
curve of CVD/cloth-Sic.
2.2
W. Liu et al.
Fig. 7. SEM micrograph of C-Sic showing between pyrocarbon and Sic.
strong
bond
little penetration. Because deposition of SIC occurs at high temperature, cooling the composite inevitably leads to microcracking the coating. Oxidation results from oxygen transport along the microcracks to the underlying C-C; in consequence, all oxidizable materials have been removed from within the sample and a hollow Sic shell has been left. The C-Sic binary matrix composites used in the present study was Sic matrix reinforced with carbon fibers separated by a carbon layer (i.e., the carbon fibers were first covered with a thin layer of pyrocarbon and then further densified by Sic). Consequently, the protective layer around each fiber consisted of an underlayer portion of pyrocarbon and overlayer portion of SIC. Owing to the small volume expansion of the carbon fiber the
Fig. 8. Typical
X-ray
diffraction
pattern
stress in the coating is reduced; hence, the cracks in the overlayer portion of Sic cannot occur. If the densified composite is machined to a desired configuration, the fiber ends will be exposed by machining. Oxidation appeared to lead to the propagation of longitudinal channels along the fiber axes and formed annular cavities. In the present study, a layer of Sic was applied to the exposed surface of the machined composites to provide a surface oxidation barrier. As a result, serious oxidation damage was not observed. From thermal cycling tests, it can be seen that C-Sic binary matrix composites show sufficient thermal shock resistance. The structural integrity of material is kept without splitting and spalling[l2]. X-ray diffractometry of samples throughout the thermal exposure testing revealed the SiOz phase present in the coating, as shown in Fig. 8. Under oxidizing conditions, at high temperature, the silicon carbide is converted to silicon dioxide, leading to the formation of a protective film of silica.
4. CONCLUSIONS
Chemical Vapor Infiltration has been successful in fabricating advanced ceramic composites. Carbonfiber-reinforced C-SIC binary matrix composites are of considerable interest because of their unique properties, such as excellent strength, sufficient thermoshock resistance, and good resistance to oxidation. A comparison of the data for different materials oxidation tests show that the mass loss rate of C-Sic binary matrix composites is lower than that of Sic-coated C-C composites. It is also noted that extensive fiber
of C-SIC binary
matrix
composites
after oxidation.
Carbon-fiber-reinforced pullout occurred in C-Sic binary matrix composites. The pyrocarbon interlayer between the carbon fiber and Sic matrix is believed to form a bond strong enough for load transfer, yet weak enough to debond readily and allow fiber pullout during crack propagation, leading to increases in strength and toughness.
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Sic composites
447
and J. J. Choury, In Proc. 3rd Int. Conf. Composite Paris (1980). 6. D. P. Stinton, A. J. Caputo, and R. A. Lowden, Am. Materials,
Ceram. Sot. Bull. 65, 347 (1986). I. E. Fitzer and R. Gadow, Am. Ceram. Sot. BUD. 65, 326 (1986). 8. D. P. Stinton, T. M. Besmann, and R. A. Lowden, Am. Ceram. Sot. BUN. 67, 350 (1988). 9. Y. M. Chiang, J. S. Haggerty, R. P. Messmer, and C. Demetry, Am. Ceram. Sot. Bull. 68, 420 (1989). 10. L. J. Schioler and J. J. Stiglich, Jr., Am. Ceram. Sot. BUN. 65, 289 (1986). 11. W. Liu, S. Sun, and M. Li, Proceedings Carbon ‘90,
Paris (1990), p. 492. 12. W. Liu, S. Sun, M. Li, and Y. Wei, Proceedings Carbon ‘92, Essen, Germany, Paper G38, Deutsche Keramische Gesellschaft (1992), p. 741.