Combustion-synthesized β′-SiAlON reinforced with SiC monofilaments

Materials Science and Engineering, A 188 (1994) 341-351

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Combustion-synthesized fl'-SiA1ON reinforced with SiC monofilaments C h a o M. H u a n g , Y o u r e n Xu, Dong Z h u a n d W a l t r a u d M. K r i v e n Department of Materials Science and Engineering and Materials Research Laboratory, University of lllinois at Urbana-Champaign, Urbana, 1L 61801 (USA) (Received November 23.1993: in revised form February 3, 1994)

Abstract Dense composites of combustion-synthesized fl'-Si3AI303N5 (fl'-SiAION) reinforced with 20 vol.% SiC monofilaments (AVCO SCS-6) were hot pressed at a temperature of 16(10 °C for 2 h under a pressure of 34 MPa. The mechanical properties of as-fabricated composites were investigated in the three-point flexure mode. The composites exhibited significant improvement in the work of fracture as well as in the ultimate strength, in comparison with monolithic fl'-SiAION. Ultimate flexure strength values between 682 and 793 MPa for the composite and between 398 and 528 MPa for the monolithic fl'-SiA1ON were obtained. A work of fracture of 13-2(I kJ m 2 was obtained for the composite, compared with 1.7-3.1 kJ m 2 for the monolithic material. Optical microscopy and scanning electron microscopy (SEM) examinations of the fractured specimens showed the usual composite toughening mechanisms of microcracking, inteffacial debonding, filament bridging and pull-out. The interracial shear strength as well as frictional stress were also investigated with a fiber push-out technique. The push-out load-deflection curves revealed a moderate interracial bonding strength of 26 MPa and a frictional sliding stress of 24 MPa. Transmission electron microscopy interfacial characterization was correlated with SEM observation of the interfacial debonding site. It revealed the presence of definite but weak physical bonding between the outermost carbon-rich layer of the SiC filament and the matrix. It appeared that the filament and the matrix were compatible with each other both physically and chemically, despite the fact that the matrix contained 20 wt.% AI203 as a secondary phase.

1. Introduction Ceramic materials based on the Si-A1-O-N system, designated sialon materials, have been of interest as a promising candidate for many special engineering applications, such as cutting tools, structural parts of internal combustion engines, and bearings [1]. Interest in this material is due to the very good inherent properties such as high strength, oxidation resistance, excellent thermal shock properties, low coefficient of friction and resistance to corrosive environments. There are three types of SiA1ON material, called fl'-, ct'-and O'-SiA1ON, based on the structures of fl-Si3N 4, a-Si3N4 and Si2N20 respectively [2]. These materials are formed as solid solutions by substitution of A1-O for Si-N simultaneously without major change in the Si3N4 or Si~N~O structure. Among these, fl'-SiAION is the most well known, fl'-SiAION is commonly described by the formula Si(, :AI=O:Ns_ =, where 0 ~
These powdered composites have poor sintering capabilities because of the free Si3N4 involved. The use of sintering aids such as yttria, ceria and magnesia allows liquid-phase sintering. These sintering aids tend to form a glassy phase which degrades the mechanical and thermal properties of the sintered body. Recently, a combustion synthesis method of preparing SiA1ON powders has been developed [3, 4]. By this method, pure single-phase fl'-SiA1ON powders can be pressurelessly sintered to almost theoretical density at relatively low temperatures ( 1550-1650 °C) without any sintering aids [4]. Concern over reliability resulting from the relative low toughness (2.5-4.8 MPa m I/2) [4] is a critical factor that limits the use of combustion-synthesized monolithic fl'-SiA1ON. Fiber reinforcement, in principle, can be employed for both strengthening and toughening a brittle matrix [5, 6]. This principle has been most successfully demonstrated in unidirectional continuousfiber composites with cement [7], glass [8] and glass-ceramic matrices [9]. Increasing attention has been directed to more refractory ceramic matrix composites, such as SiC-monofilament-reinforced Si3N 4 composites. The composite of reaction-bonded Si3N 4 (RBSN) reinforced with SiC filaments has been shown © 1994 - Elsevier Sequoia. All rights reserved

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to yield a material which possesses greater strength and toughness than unreinforced RBSN of the same density [10]. However, in the as-fabricated state, the matrix has a porosity level of about 30 vol.%, which results in poor oxidation resistance and low matrix strength [ 11 ]. Hot-pressed Si3N 4 reinforced with SiC filaments exhibits a significantly higher work of fracture (WOF), but the strength is moderate in relation to typical strengths achieved in hot-pressed Si3N4 [12-14]. High temperatures (1750°C or above) are generally required to make a fully dense Si3N4 matrix, from which the strength of filaments is degraded, and thus the fracture strength of the composite is limited [13]. In addition, the use of sintering aids has limited the composite high temperature capabilities. /3'-SiAION exhibits most of the mechanical and thermal properties of Si3N4, but superior sinterability to Si3N4 [1, 2]. This advantage makes/3'-SiAION more attractive than Si3N 4 as a matrix material in a composite. A composite system is generally chosen on the basis of the commercial availability as well as compatibility requirements for fibers and matrix. In general, high modulus, high strength and high temperature stabilities are required for the fiber. On the contrary, high temperature capabilities including creep and oxidation resistance, as well as a moderate thermal expansion difference relative to the fiber are required for the matrix. The composite/3'-SiA1ON reinforced with SiC monofilaments meets most of these requirements, and potentially it will be superior to the composite of Si3N4 reinforced with SiC monofilaments. However, research efforts on this composite system have been relatively few to date. No work on the composite of combustionsynthesized fl'-SiA1ON reinforced with SiC monofilaments has yet been published. The present paper reports on the results of an investigation in which SiC monofilaments were used to reinforce a combustion-synthesized fl'-SiA1ON (Si3A1303Ns) matrix. The objective of this study was to demonstrate the feasibility of such a composite system. Mechanical properties, interracial bonding strength, frictional sliding stress and interfacial microstructure were examined in a hot-pressed/3'-SiAION (Si3AI303_ Ns) composite reinforced with (AVCO SCS-6) SiC monofilaments.

2. Experimentalprocedure Commercial-grade SiC monofilaments (AVCO SCS6) were obtained from Textron Specialty Materials, Lowell, MA. These filaments were made by a chemical vapor deposition process in which SiC about 50/~m thick was deposited onto a carbon core of 37 /~m diameter, followed by depositions of carbon and

carbon-silicon layers about 3/~m thick resulting in an overall filament diameter of 140 pm. A /3'-SiAION (Si3A1303Ns) powder (Benchmark S-011) was obtained from Benchmark Structural Ceramics Corporation, Buffalo, NY. The powder was made primarily by a combustion process involving an aluminothermic reduction of silica followed by in-situ synthesis of/3'SiAION. The as-supplied powder contained 80 wt.% /3'-SiA1ON (Si3A1303Ns) and 20 wt.% A1203, as reported by the manufacturer. The" SiC continuous filaments were washed in trichloroethylene to remove any trace of organic impurities and subsequently washed alternately in nitric acid (0.01 N) and ammonium hydroxide (0.1 N) to remove inorganic trace impurities. The final washing was done with distilled water. The surface-cleaned fibers were dried in a vacuum oven at 160°C overnight. The cleaned filaments were cut into short 2 in filaments, which were unidirectionally arranged in one plane and then stacked alternately with matrix powders in a square (2 in x 2 in) graphite die. Hot pressing was carried out at a temperature of 1600 °C for 2 h under a pressure of 34 MPa in a nitrogen atmosphere. The filament content of the composite was about 20 vol.% for all the samples. For comparison, monolithic fl'-SiAION was also hot pressed in the same manner without filaments in this study. The densities of the as-fabricated monolithic material and the composites were determined by the Archimedes method. Mechanical testing of the uniaxially reinforced composite was performed in the three-point flexure mode. Hot-pressed composite plates were surface ground to a finish of 60/~m and then cut into individual bars parallel to the filament axis. The edges of the tensile face of each specimen were slightly bevelled on a 9 /zm grit diamond lap. Typical dimensions of bar specimens were 50 mm long, 4.8 mm wide and 2.8 mm thick. The span of the lower support pins was 30 mm, which gave a span-to-thickness ratio of 10.7 in the three-point flexure mode. The flexure tests were performed in a universal testing machine (model 4502, Instron Corporation, Canton, MA) at a cross-head speed of 0.0127 cm min i (0.005 in min ~). Load-deflection data were obtained at room temperature for monolithic /3'-SiAION and the composite samples which were processed under similar conditions. A number of mechanical properties were determined from the load-deflection data. These included the stress of first matrix cracking, the ultimate stress and the WOE The stresses were calculated from the load, which was measured with a load cell, based on the simple beam theory. The WOF was obtained by dividing the area under the load-deflection curve by the cross-sectional area of the sample. The mode of failure (tensile or shear) was determined from visual observation during

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the test and optical microscopy examinations of the failed samples subsequent to the test. Scanning electron microscopy (SEM) examination was also conducted on as-fractured and polished cross-sections of failed composites to determine the nature of the fracture process, and the toughening mechanism. The interracial properties, such as shear strength and frictional stress were measured by the single-fiber push-out technique, which has been described elsewhere [ 15-17]. The measurements were made in a universal testing machine (model 4502, lnstron Corporation, Canton, MA), with a flat-end probe, l(t0 um in diameter and made of tungsten carbide. In this study, for the purpose of in-situ alignment of probe and filament, a simple modification was made based on the design of Bright et al. [17]. In this, two two-axis micropositioners were used instead of one. One two-axis micropositioner, which carried a stereomicroscope and travelled on two parallel feed shafts, was used for the filament and probe alignment purpose. The other, on which the sample was mounted, was used for the filament positioning purpose. Specimens for push-out tests with a thickness of 0.5 mm were prepared by cutting thin sections perpendicular to the filament direction in the uniaxial composites. The as-cut sections were surface ground on both sides, and polished to a final finish of 1 /~m. The push-out specimen was mounted on a slotted alumina holder which was centered on one of the two-axis micropositioners. The push-out tests were conducted at a constant cross-head displacement rate of 60 /,m rain ~ in an ambient atmosphere. Typically, seven to eight filaments were pushed out in order to obtain average properties for

Fig. 1. Scanning electron micrograph of a consolidated SiC filament, combustion-synthesizedfl'-SiAION composite (polished cross-section).

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the composite. The interracial shear stress was calculated from the load required to initiate filament motion and the circumferential area of the filament in contact with the matrix. The sample was then turned over and measurements were repeated on filaments which had already been pushed out, from which the frictional stress could be evaluated. After push-out testing, a scanning electron microscope (H-800, Hitachi Ltd., Japan) was used to examine the debonded interface and to locate the origin of debonding. High resolution transmission electron microscopy (HREM)(CM-12, Philips Electronic Instruments, Eindhoven, Netherlands) examination was also conducted to determine the microstructure of the interface between the SiC filaments and matrix. The specimen for HREM examination was prepared in the conventional manner.

3. Results and discussion 3.1. Physical prol)erties A scanning electron micrograph showing the crosssection normal to the filaments of a dense consolidated composite is given in Fig. 1. It has a random distribution of filaments in a dense fl'-SiA1ON matrix, and the filaments are completely isolated by the matrix. Filament-to-filament contact, which can degrade composite properties, was not found through the examination of cross-sections of a number of specimens. The composite samples had a density of 3.18 g cm ) which was approximately 97% of the predicted full density of 3.26 g c m ~, based on a ff-SiAION density of 3.28 g cm 3 and a silicon carbide filament density of 3.2 g c m 3. Variations in specimen-tospecimen density were small, which was indicative of uniform consolidation of the composite. The measured density of about 3.2 g cm 3 for the monolithic fl'SiAION was close to its theoretical density. The X-ray diffraction analysis from as-fabricated monolithic samples showed fl'-SiAION as the major phase plus 15-20 wt.% AleO 3 as a minor phase. A summary of the mechanical properties of both monolithic fl'-SiAION and its composites is given in Table 1. The monolithic ff-SiAION fractured at strength values between 398 and 528 MPa, while the SiC monofilament-fl'-SiA1ON composites failed at ultimate strengths between 682 and 793 MPa. The toughness of the composite was significantly higher than that of monolithic fl'-SiAION, which was reflected by the measurement of the WOE WOE values for the composite were measured to be 13-20 kJ m -+, compared with 1.7-3.1 kJ m -" for monolithic material. The mean stress of first matrix cracking in the composite was 346 MPa, which was relatively lower than the fracture strength of the monolithic materials (1475

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fl'-SisAl~O¢A(~reinforced with SiC monofilaments

TABLE 1. Mechanicalproperties of SiC-monofilamentreinforced fl'-SiAIONcompositeand monolithicmaterial Material

Specimen Crack Ultimate W O F initiation strength ( k J m - 2 ) strength (MPa) (MPa)

800

.

.

.

,

.

- -

.

.

.

.

,

.

.

~

.

,

.

.

.

,

.

.

.

.

.

,

.

.

.

.

,

.

.

.

.

,

.

.

.

.

SiC / B'-SiA1ON Composite

600

g 400

Composite

C-01 C-02 C-03 C-04 C-05 C-06 C-07 C-08

Mean value Standard deviation Monolith

Mean value Standard deviation

M-01 M-02 M-03 M-04 M-05 M-06 M-07 M-08

317.8 337.1 340.4 440.0 346.2 277.2 330.0 382.7

687.0 741.4 682.3 762.9 775.6 780.8 793.8 758.2

15.7 16.4 12.1 18.2 20.5 17.2 19.5 17.0

346.4 16.9

743.8 41.0

16.7 2.7

439.6 528.7 497.8 459.4 534.6 452.9 490.3 398.7

2.21 3.55 2.73 2.37 3.19 2.34 2.63 1.80

475.3 46.3

2.67 0.57

MPa). The fracture surfaces of a selected number of bend test specimens were examined in an optical microscope to seek an explanation for the reduced strength of the matrix. In all the specimens examined, fracture appeared to initiate at the site of grinding damage to the filaments or filament troughs on the surface. The finished surface of composite specimens for flexure testing was particularly rough compared with the monolithic specimens, because the SiC filaments tended to be pulled out during grinding. Therefore the rough surface finish of the composite specimens was believed to be responsible for the low strength of the first matrix cracking in the composite. 3.2. Fracture behavior The general features of the load-deflection curves for three-point flexure tests for monolithic fl'-SiAION and a SiC monofilament-fl'-SiAION composite are shown in Fig. 2. For the monolithic material, all load-deflection data showed elastic behavior up to the point of fracture, which is typical of catastrophic fracture behavior for brittle ceramics. For the composite, as a comparison, an initial linear elastic region was followed by a non-linear load increase to a maximum, followed by a continuous load decrease. The noncatastrophic decrease in load gave the composite the appearance of being "tough". The linear elastic behav-

200

0.0

0.1

0.2

0.3

0.4

0.5

Displacement

(ram)

0.6

0.7

0.8

Fig. 2. Load-deflection behavior in three-point flexure for the monolithic fl'-SiAIONand SiC-fl'-SiAION and SiC-/%SiAION composites.

ior occurred up to a small but distinct discontinuity, where the load dropped slightly. This discontinuity was a clear indication of the first matrix cracking. Beyond this discontinuity, the load-deflection curve exhibited a saw-tooth-type non-linear behavior, accompanied by multiple matrix cracking and interfacial debonding events, before a sudden drop-off in the load. The sudden drop in load was associated with unstable matrix cracking and some filament fracture. Considerable additional inelastic deformation beyond the maximum load was seen for most specimens, which could be attributed to pull-out of broken filaments. The examination of failed specimens provided an insight into the characteristics of the load-deflection curve and thus toughening mechanisms. Visual observation of samples during the flexure test revealed that almost all the specimens fractured in a tensile mode under the middle loading rather than on the compressive side via buckling and shear delamination. This was confirmed by optical microscopy examinations. Buckling and delamination mode failures of composite samples tested in three-point flexure have been observed by several previous investigators [18, 19], and a large span-to-height ratio has been suggested to avoid this undesired fracture mode. It is generally agreed that testing in pure tension is a proper way for determination of ultimate strength of reinforced composites. Optical micrographs (edge view) of the failed composites tested in three-point flexure are shown in Fig. 3. These micrographs were taken on the specimens which were subjected to deformation considerably beyond the maximum load. The micrographs clearly indicate that, for most specimens, the major crack was initiated on the tensile surface under the middle loading point and was arrested at the midplane

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fl'-Si3Al~O¢N ~reinforced with SiC monofilaments

r[

2mm

Ittlt it[lltii llbltLtlL1lJtLiIllitllil lllll tlliiilllLItllllllttqitltlllllitllltlltLil ;llll I zo 31o 4 o c., 50 t 6to ,,,,21o .... 310 ....... 41o.... 51o, 61o Fig. 3. Optical micrographs of failed composites in the edge view.

of the beam. Another observation by optical microscopy of the tensile surface (Fig. 4) was the formation of a series of parallel microcracks away from the primary crack in the central test area. These parallel microcracks were more or less equally spaced, and their average distance was measured to be 2 mm. They became progressively narrower in width with increasing distance from the primary crack. In addition, these microcracks penetrated only to the neutral plane of the beam. Figure 5 is a scanning electron micrograph of the primary crack on the tensile surface of a fractured composite specimen. A wide open primary crack formed across the width of the specimen which was bridged by the filaments, accompanied by individual filament fracture. An as-fractured cross-section surface of the composite is shown in Fig. 6. Filament pull-out, clean debonded filament-matrix interfaces and a stepped fractured surface are apparent. The non-planar fracture surface was a clear indication of crack deflection by the reinforcing filaments during the crack propagation. The pulled-out filaments retained their original nodular surface morphology (Fig. 7(a)), which was mirrored by the morphology of the trough remaining in the matrix (Fig. 7(b)). This suggested that the debonding was located at the site between the carbon layer of the filament and the matrix. The intact morphology of pulled-out filaments was an indication that the filaments had been well protected from the damage during the processing. Figure 8(a) is a scanning electron micrograph of a polished cross-section of the failed composite, showing the processes of crack propagation and deflection. The cracks always propagated from one filament to the next and were diverted by the filaments. All the propagating cracks were deflected

m

Fig. 4. Optical micrograph showing periodic matrix microcracking (arrowed) normal to the filament axis.

Fig. 5. Scanning electron micrograph showing filament bridging in the failed composite.

around the filaments, and there were no locations where cracks had gone into the reinforcements. Higher magnification SEM (Fig. 8(b)) indicates that the crack deflected along the interface between the carbon layer of the filament and the matrix, and this was consistent with the observation of as-fractured surfaces. Based on these observations of fracture behavior, the progressive nature of SiC monofilament-reinforced fl'-SiA1ON matrix composite failure, which was reflected by the load-deflection curve, was contributed to by microcracking, crack deflection, interfacial debonding, filament bridging and pull-out. The filament bridging and pull-out were obviously the dominant toughening mechanisms in this composite.

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Fig. 6. Scanningelectron micrograph of a failed cross section (as fractured) of SiC-fl'-SiAION, showing filament pull-out and stepped matrix surfaces.

However, crack deflection and microcrack processes also contributed to the toughening to a certain extent. 3.3. Interfacial characterization Figure 9 shows a typical load-deflection record obtained during a push-out and a push-back test on a single-filament in a specimen with 0.5 mm thickness. The curves exhibited the general features of a push-out and push-back test. The push-out curve consisted a linear elastic loading of a bonded filament and a nonlinear loading of a filament sliding against the matrix. The linear stage terminated at a load Pa where the load dropped catastrophically to a lower load. This load drop corresponded to complete debonding between the filament and matrix. Following the debonding event, the load rose slightly to a peak P~, which corresponded to the beginning of relative sliding between the filament and the matrix. Beyond this point the load decreased as the sliding filament exited from the bottom side of the specimen. The distinct separation of debonding and sliding peaks, as shown in Fig. 9, indicated the presence of definite interfacial bonding between the filament and matrix. Upon push-back, the load increased steadily until the filament reached the

Fig. 7. Scanning electron micrographs showing (a) morphology of pulled-out SiC filamentsand (b) mirror image of the filament surface seen in the matrix trough above the other half of the broken filament.

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original location. At that location, there was a load drop associated with the fiber resetting into its initial position, as was observed in previous studies [20, 21]. The load then rose to a peak Pt, and decreased gradually as the filament slid out of the specimen. The period and magnitude of the seating drop have been correlated with the interracial roughness and amplitude of the roughness respectively by previous investigators

[20,21]. The values measured for the seating drop periodicity were narrowly dispersed and had a mean value of 9.5/~m. However, the magnitudes of the seating drops were widely scattered and varied from 1.5 to 3.5 N. A summary of interracial properties of composites fabricated in this study is given in Table 2. The first load Pa obtained from the push-out data, was used to calculate the interfacial shear strength rd, based on the formula

Pd 6 5

Ps

- -

~

Push-out Push-back

Period = 10 gm

1

0.00

0.05

0.10

0.I5

0.20

Displacement (mm) Fig. 9. Typical load-deflection behavior of push-out and corrc'sponding push-back of an SiC filament in the fl'-SiAION matrix.

T A B L E 2. lnterfacial shear strength and frictional stress of an

i¸¸

SiC-monofilament-reinforced fl'-SiA1ON composite Filament

Fig. 8. Scanning electron micrograph of a failed polished crosssection of an S i C - I f - S i M O N compositc, showing (a) crack propagation and deflection and (b) the site of interracial debond-

ing.

rL~ (MPa)

rs

rb

(MPa)

(MPa)

1 2 3 4 5 6 7

28.6 29.3 27.3 24.2 22.5 24.3 31.1

24.1 25.0 24.8 23.8 22.5 23.(I 28.(I

17.3 18.5 18.2 16.4 16.3 16.3 2(I.3

Mean value Standard deviation

26.5 2.8

24.5 1.7

18.1 0.6

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fl'-Si3Al30~Ns reinforced with SiC monofilaments

Fig. 11. Scanning electron micrograph showing the surface morphologyof a pushed-out SiC filamentand visiblewear tracks (arrowed).

r d-

I'd 2zrRt

where R is the filament radius, and t is the sample thickness. This approach assumed a uniform shear across the entire interface when the load was applied, and was limited to a thin specimen with a fiber-lengthto-fiber-diameter ratio t / 2 R smaller than 4 [22]. The specimen used in this study had a length-to-diameter ratio of 3.6, which fulfilled this approximation. The mean value of interfacial shear strength calculated was 26 MPa. As a comparison, a simple method derived by Aveston et al. [23] was also used to calculate the interfacial shear strength, in which Vm ~rmR Vf 2L

Fig. 10. Scanning electron micrograph of SiC filaments after push-out, showingthe site of debondingand sliding:(a) top view; (b) bottom view.

where o mis the ultimate strength of the matrix, R is the filament radius, L is the mean separation of microcracks and Vm and Vf are the matrix and filament volume fractions respectively. After substituting the appropriate values in this formula, the calculated value of r d was found to be 24 MPa, which was very close to the value (26 MPa) obtained from the fiber push-out data. The frictional sliding stress rb obtained from the push-back Pb was compared with r~ obtained from the push-out P~. The rb value of 18 MPa was a little smaller than the r~ value of 24 MPa, which suggested the existence of interfacial abrasion during filament sliding. SEM examination (Fig. 10) shows that the site of debonding and sliding was located at the interface between the carbon surface layer of the filament and the matrix, which was consistent with the observation

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fl'-Si~A(~O~N~reinforced with SiC monofilaments

of filament pull-out in the flexure test. The pushed-out filament retained its original nodular surface morphology (Fig. 11 ), but with visible abrasion traces (arrowed in Fig. 11 ). A maximum nodule size of 10 /~m was measured, comparing well with the 9.5 ktm periodicity of the seating drop. The mirror images of pushed-out filament and trough surfaces were evidence for weak bonding between the filament and matrix. The interface between If-SiMON and SiC filament was further examined by HREM. A lattice image of a near-interfacial region of a If-SiMON grain and the outermost turbostratic carbon (TC) layer of the SiC filament is shown in Fig. 12. The electron beam was parallel to a (001) zone axis of the If-SiMON and the lattice fringes of(100) planes were parallel to the interface. The interface between the fl'-SiAION grain and the TC layer of the SiC filament was quite well defined,

349

as shown in Fig. 12(a). Obviously, there was no extended reaction zone between them. A similar lattice image of the interface between the AI203 grain and TC layer of a SiC filament was also obtained (Fig. 12(b)), in which the lattice fringes of (003) planes were parallel to the interface. The morphology of the carbonaceous layer of SCS-6 SiC filament was similar to previous HREM observations [24, 25]. The outermost layer of the filament was characterized as turbostratic carbon with basic structural unit blocks constituting a heavily textured appearance. TC can take on forms ranging continuously from a near-amorphous to a highly crystalline state [25], which explains the partial presence of amorphous phase in the TC layer as was observed in HREM images. The continuity across the matrix and filament, as shown in the HREM images, indicated the presence of definite interracial bonding between SiC

Fig. 12. HREM images of interfacial regions of the SiC-fl'-SiAION composite: (a) interface between a fl'-SiAION grain and TC; (b) interface between a minor-phase Al203 grain and the TC.

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monofilaments and the matrix, which was consistent with the observation in the push-out test. However, apparent lack of a chemical reaction layer between the filament and matrix suggested that the bonding between them was simple physical adhesion formed mainly by consolidation of the composite during hot pressing. The matrix seemed to be compatible with the carbon-coated SiC monofilament, despite the fact that the matrix was a mixture of ff-SiAION (80 wt.%) and AI203 (20 wt.%). The chemical compatibility between the Al203 and SiC has been demonstrated in SiCwhisker-reinforced A1203 composites [26, 27]. Unlike A1203-SiC, ff-SiA1ON and SiC forms a thermodynamically unstable system when sintered in an N 2 atmosphere [28]; that is, when combined at high temperature, a chemical reaction occurs between SiC and the matrix. By introducing a carbon layer between them, the reaction can be inhibited. It has been reported [29] that a composite of fl'-SiAION reinforced with 30 vol.% uncoated SiC whiskers exhibited slightly less toughness than the monolithic fl'-SiAION. Strong bonding due to the chemical reaction between the SiC whiskers and matrix, and inhibition of /3'SiAION grain growth were considered to be the main reasons for this. fl'-SiA1ON composites with the same amount of carbon coated whiskers, however, had a significantly increased fracture toughness. The improved fracture toughness of carbon-coated whisker-reinforced composites was found to be due to slippage or flaking in the carbon interface between the whiskers and the matrix. In the case of the present study, the carbon layer of the SiC filament was considered to play a key role in improving the toughness of fl'-SiA1ON composites, which prevented any chemical reaction between the SiC and the fl'-SiAION matrix, thereby providing weak interfacial bonding and enhancing debonding and filament pull-out. The secondary phase of Al203 in the matrix might have some impact on the frictional stress of filament sliding. As a result of the larger coefficient of thermal expansion ( C T E ) ( ( 8 - 9 ) × 10 -6 °C -1) of A1203, the mixture of ff-SiA1ON and Al203 was expected to have a higher mean CTE value than pure fl'-SiA1ON (3.04× 10 -6 °C-I). A larger matrix CTE generally results in a higher residual compressive stress on a fiber and correspondingly a higher frictional stress would be expected. This was reflected by the fact that the composite made in this study had a much higher frictional stress compared with the hot-pressed composites of Si3N4 reinforced with SiC monofilaments [14], where the CTE value of Si3Na ((3.1-3.2)x 10 6 oc-z) was very close to ff-SiAION. The filament sliding stress has a large effect on the fracture toughness because it determines the maximum load which the sample could

bear at fracture. It was reported [14] that the fracture toughness increased with increasing frictional stress in the case of the composite which experienced filament pull-out at fracture but, when the frictional stress became too large for the filament to be pulled-out, the fracture toughness of composites was not affected by the filament-matrix interfacial frictional stress. Therefore a certain amount of secondary-phase A1203 contained in the as-received fl'-SiAION matrix might have some positive effects on the composite toughening. Further confirmation of this postulation will be given in a separate paper.

4. Conclusions

A combustion-synthesized fl'-SiAION (Si3AI303Ns), which contained 20 wt.% A1203, was uniaxially reinforced with SiC filaments and hot pressed into a dense composite. The high sinterability of the /3'SiAION allowed the composite to consolidate to nearly full density at a relatively low temperature (1600 °C or less), thereby protecting SiC filaments from degradation. Three-point flexure testing was conducted to determine the mechanical properties. Optical microscopy of failed samples subsequent to the test illustrated that the failure occurred via the tensile mode below the middle loading pin. The composite exhibited a significantly improved fracture toughness, as reflected by the WOF, and improved ultimate fracture strength, but relatively lower crack initiation resistance, compared with the monolithic ff-SiA1ON materials. This was found to be related to the rough surface finish of the specimen. Examination of failed samples showed that toughening in the composite involved a combination of filament pull-out, crack deflection and possibly microcracking mechanisms. Interfacial debonding and filament sliding against the matrix were clearly distinguished in push-out load-displacement curves. Correlation of TEM observations with push-out load-displacement curves revealed the presence of definite but weak filament-matrix bonding in the asfabricated composite. The lack of a chemical reaction layer between the filament and matrix suggested that the bonding was likely to be physical adhesion caused during consolidation of the composite. SEM examination indicated that debonding either in push-out tests or in flexure tests was located at the interface between the outermost carbon layer of the filament and the matrix, which further confirmed the presence of weak interfacial adhesion. A secondary-phase A1203 contained in the matrix appeared to upgrade the frictional stress of filament sliding, which was thought to have a positive effect on the composite toughening.

C. M. Huang et al.

/

fl'-Si3A,'~O.,,N, reit{forced with SiC monofilaments

Acknowledgments F u n d i n g for this r e s e a r c h was p r o v i d e d b y the U S A i r F o r c e o f Scientific R e s e a r c h t h r o u g h G r a n t AFOSR-F49620-93-1-0227. T h e a u t h o r s wish to t h a n k Dr. A . K u m n i c k f r o m T e x t r o n S p e c i a l t y M a t e r i a l s f o r s u p p l y i n g t h e SiC filaments. T h e a u t h o r s w o u l d also like to t h a n k D. H. K u o a n d J. L. Shull for t h e i r a s s i s t a n c e in p r e p a r i n g t h e c o m p o s i t e s p e c i m e n s a n d o p t i c a l m i c r o s c o p y o b s e r v a t i o n respectively. V a l u a b l e d i s c u s s i o n s with P r o f e s s o r T. J. M a c k i n f r o m t h e D e p a r t m e n t o f M e c h a n i c a l E n g i n e e r i n g at U n i versity o f Illinois at U r b a n a - C h a m p a i g n a r e a p p r e ciated. U s e of the e l e c t r o n m i c r o s c o p y facilities at the C e n t e r for M i c r o a n a l y s i s o f M a t e r i a l s in the M a t e r i a l s R e s e a r c h L a b o r a t o r y at U n i v e r s i t y of Illinois at U r b a n a Champaign are acknowledged.

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