Reaction-infiltrated, net-shape SiC composites

MATERIALS SCIENCE & ENGINEERING ELSEVIER

Materials Science and EngineeringA195 (1995) 131-143

A

Reaction-infiltrated, net-shape SiC composites Leszek Hozer, Jonq-Ren Lee, Yet-Ming Chiang Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, MA 02139, USA

Abstract

The processing and net-shape capability of SiC composites formed by reactive infiltration of glassy carbon preforms using liquid silicon alloyshave been studied. Compositesconsisting of 65-85 vol.% SiC with a secondary phase of Si, Si-A1, or Si-Cu compounds have been prepared. Tests of dimensionalchanges and the retention of surface finish during reactive infiltrationhave been conducted. Changes in linear dimensions between a machined preform and the reacted composite can be smaller than 0.1%. Surface finishes Ra< 1.8/~m are achievablein as-reacted composites. The simultaneousjoining and reaction bonding of parts to form a more complex shape has been demonstrated. The flexural strength, modulus of elasticity, hardness, and indentation fracture toughness of these materials have been characterized.

Keywords: Siliconcarbide; Composites;Infiltration

I. Introduction

Silicon carbide composites have a high potential as advanced ceramics for demanding applications. Their superior mechanical properties at low and high temperatures, wear resistance, thermal properties, and corrosion resistance have attracted the attention of many researchers. However, the processing of structural silicon carbide is complex and costly. Pressureless sintering requires high temperatures (over 1800 °C) in reducing atmospheres. To obtain high density material, pressure sintering is also frequently used, which, however, increases the cost and limits the dimensions and shape of the final product. Conventionally manufactured parts frequently require further machining, typically using diamond tooling. The high processing temperatures and lack of net-shape capability in conventional processing are primary concerns confining the number of low cost applications of silicon carbide. Fabrication of reaction-bonded silicon carbide (RBSC) by infiltration techniques is a lower temperature route which has been extensively studied in the past [1-6]. The majority of studies have been concerned with so-called siliconized silicon carbide pro0921-5093/95/$9.50 © 1995 - ElsevierScienceS.A.All rightsreserved SSD10921-5093(94)06512-8

duced by reacting liquid silicon with a carbon-containing preform. The reaction of carbon with liquid silicon is very fast; studies by Ness and Page [7], Pampuch et al. [8], and recently Chiang et al. [6] support a mechanism based on solution and reprecipitation. In a variant of the reaction bondh~g process Hucke [4,5] also used polymer-derived microporous carbon preforms and achieved SiC-Si composites with flexural strengths of 600-800 MPa at room temperature, superior to that of most RBSCs (about 300-400 MPa). SiC-Si material is of interest when, for example, corrision, strength, hardness, and wear resistance of the product are of importance whereas high fracture toughness is not critical. Examples of such products can be nozzles for corrosive agents or high temperature gases, combustion gas heat exchange tubes [9], or seals. Messner and Chiang [10,11] and Chiang et al. [6] have also demonstrated the use of alloyed melts in order to control the constitution of the secondary phase in RBSC. In the silicon-molybdenum-carbon system [6,10,11] it was shown that, when porous carbon is infiltrated with Si-Mo melts, silicon carbide is formed as a primary phase while the alloying element is rejected into the remaining melt, forming a second-

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ary phase enriched in the additive, eventually forming MoSi 2 as a refractory secondary phase in an SiC composite. It is expected that the introduction of an interconnected metallic phase into RBSC, in place of Si, would improve its fracture toughness. In this work we explore the use of Si-AI and Si-Cu alloys for reactive infiltration, in order to form composites with AI- and Cu-rich secondary phases. A n additional role of the alloying is to control wetting and infiltration by the liquid. Most liquid metals exhibit a high contact angle with carbon, therefore requiring the application of pressure to achieve infiltration [12]. Silicon, because of its reactivity with carbon, is an exception which spontaneously infiltrates porous carbon. While pure aliminum and copper do not, it is found that, by alloying these metals with silicon, spontaneous infiltration can be achieved. The specific phase combination of SiC-A1 has also been widely studied as a metal matrix composite, in which the SiC phase is usually discontinuous. The materials discussed in this work have, on the contrary, an interconnected structure of both the SiC and the secondary phase, which provides properties more characteristic of a cermet. The SiC-A1 combination is also a promising material for substrates for electronic devices [13]. In this application, the high thermal conductivity is of interest, as well as the good thermal expansion coefficient match to silicon or GaAs. Aluminum has a high thermal expansion coefficient relative to that of SiC, Si, or GaAs, and the controlled introduction of aluminum into silicon carbide enables systematic variation of the expansion coefficient of the composite [14]. In this paper we discuss the processing, microstructure, phase constitution, and properties of SiC composites produced by reactive infiltration. In addition, we have characterized in detail the dimensional retention and surface finish due to processing. The netshape processing of prototype components, and the simultaneous joining and reaction bonding of complexshape parts, is demonstrated.

2. Composite processing 2.1. Preform preparation

The fabrication process of the composites is shown schematically in Fig. 1. The microporous carbon preforms used in this work were prepared after a process originated by Hucke [4], discussed in Ref. [6], and therefore only briefly summarized here. Liquid mixtures of furfuryl resin and furfuryl alcohol (furfuryl resin (QUACORR ® 1300) and furfuryl alcohol (QO FA®) from QO Chemicals, Inc., West

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Lafayette, IN) were used as one component of the preforms; this constituent forms the glassy-carbon skeleton after pyrolysis. The furfuryl solutions are mixed with solutions of ethylene glycols, which are initially miscible with the furfuryl but become immiscible on polymerization of the furfuryl at quite low temperatures. We have directly observed, using optical microscopy, that miscible solutions of furfuryl resin, furfuryl alcohol, and ethylene glycols undergo phase separation at temperatures lower than 100 °C. Concurrently, the viscosity of the furfuryl-rich phase increases, eventually forming a rigid skeleton. The glycol-rich phase remains fluid and volatile, and can be readily removed from the open pore structure. On pyrolysis the furfuryl-rich phase transforms to glassy carbon, forming a microporous carbon reticulated structure (see Fig. 3). It is possible to vary the morphology of the microstructure, the overall preform density, and the pore size by varying the ratios of the components and the temperature cycle during polymerization. The preforms were prepared in simple cylindrical shapes by casting the initial liquid mixture of furfuryl and glycol into molds. A subsequent temperature cycle of polymerization and pyrolysis was carefully designed on the basis of thermogravimetric analysis (TGA) measurements, microstructure observations, and measurements of density and porosity. The results of a

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typical T G A analysis are shown in Fig. 2. Since substantial weight loss is o b s e r v e d at t e m p e r a t u r e s below 150 °C, where volatile liquid products are removed, a slow heating rate through or dwell steps in this temperature range are necessary to avoid m a c r o s c o p i c flaws in the preform. A cast p r e f o r m is first polymerized at 40 °C for 4 h. T h e t e m p e r a t u r e is then raised to 70 °C and held for the next 12 h to complete the polymerization and phase separation. Scanning electron m i c r o s c o p y (SEM) observations show that at this stage the phaseseparated skeleton is fully formed, although liquidfilled pores remain. T h e rigid but unpyrolyzed p r e f o r m is r e m o v e d f r o m the mold, and the bulk of the liquid glycol phase is extracted f r o m the o p e n porosity with absorbent media. Residual glycol phase is volatilized during the subsequent pyrolysis, which is p e r f o r m e d in an N 2 a t m o s p h e r e up to a m a x i m u m t e m p e r a t u r e of 1000 °C in a resistance-heated box furnace equipped with a steel retort. T h e complete t e m p e r a t u r e cycle lasts about 48 h. We found that a final annealing in v a c u u m (5 × 10 -3 Torr) at 1 6 0 0 - 1 7 0 0 °C for about 1 h was useful to r e m o v e gas-forming species which otherwise interfere with infiltration. T h e total weight loss during the thermal cycle of pyrolysis is on average equal to about 70%, and the p r e f o r m shrinks by about 3 0 % f r o m the mold dimensions. T h e cylindrical p r e f o r m s p r e p a r e d u n d e r these laboratory conditions had m a x i m u m dimensions of 15 m m diameter and 100 m m length. A variety of examinations showed that they were uniform in density and microstructure. T h e microstructure of the pre-

Fig. 3. Examples of various achievable microstructures of the porous glassy carbon preforms: (a) spherical particulates; (b) interconnected network; (c) interconnected network with a secondary closed porosity not accessible for infiltration.

forms can be varied f r o m an interconnected carbon network to spherical particulates, as shown in Fig. 3. For the purposes of this work, we chose the type of

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spherical particulate microstructure shown in Fig. 3, and varied the preform densities from 0.80 to 0.73 to 0.65 g cm -3. The corresponding average pore size (measured using mercury intrusion porosimetry) ranged from 1 to 3/~m. The macroscopic density of the preforms could be reproduced to _+0.02 g cm -3 between different batches. The specific carbon density of the preform skeleton was found to be approximately 1.5 g cm -3, using mercury porosimetry. The pyrolyzed preforms, while brittle, could be machined using tooling typical of that used for the machining of graphite. In Fig. 4 we demonstrate that fairly complicated shapes including threaded parts are readily machined. As discussed later very good dimensional accuracy and surface finish can be achieved after the reaction process. This is of great importance from the point of view of net-shape capabilities of the reactive infiltration approach.

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2.2. Alloy composition systems

In the absence of any carbon loss or dimensional change during reaction, the overall carbon density of the final composite is simply that of the starting preform. Thus the carbon density is one boundary condition for determining the composition of the infiltrating alloy. In the case of SiC-metal composites our goal was to introduce a metal-rich secondary phase into the SiC interpenetrating network. If the reaction proceeds to completion, the compositions and amounts of the individual phases are determined by the overall composition of the composite and the applicable phase equilibria at the reaction temperature. Thus we targeted overall compositions in the Si-A1-C and Si-Cu-C systems which, based on available phase equilibria, were likely to give two-phase mixtures of SiC and a metal-rich liquid at the processing temperatures (i.e. they lie in two-phase fields). The carbon densities and corresponding alloy compositions are shown in Table 1. The final phase assemblage after cooling is also determined in part by the composition and solidification process of the liquid, as discussed later. In the Si-AI-C and Si-Cu-C systems, no ternary compounds were expected, and none were observed in the final composites. The assumption that all the carbon is reacted and converted to SiC was supported by examination of the reacted composites. Several samples were made at each of the compositions shown in Table 1, and representative results will be discussed below. 2.3. Reactive infiltration

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Machined preforms were mounted on a graphite holder and reaction infiltrated by dipping into the melt, contained in a boron nitride crucible in a vacuum furnace at 10-2-10 -3 Torr (Thermal Technology, Astro Division, Santa Barbara, CA). The infiltration temperature and time are dictated by two issues. First, the temperature must be sufficiently high to have a single-phase liquid and to allow spontaneous infiltration. Second, the temperature and time of infiltration must be sufficient for substantially complete reaction. In the case of pure silicon, a temperature of 1450 °C, just above the silicon melting point (1410 °C), and reaction times as short as 5 min gave a completely infiltrated and reacted SiC-Si. In the case of the metal-silicon alloys, Si-A1 and Si-Cu, the liquidus temperature is relatively low for the compositions in Table 1 (1100-1325 °C) [15], but spontaneous infiltration required superheating above the liquidus temperature in most cases. While a complete study of the influence of composition and temperature on wetting and infiltration in these two alloy

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Table 1 Melt compositionsand infiltrationconditionsused for metal-siliconalloyinfiltratedsamples Sample

Melt composition (mol.%)

Preform density (gem -3)

Infiltration temperature (°C)

Total infiltration and reaction time (min)

SiC-Si SiC-All SiC-AI2 SiC-AI3 SiC-AI4 SiC-Cul SiC-Cu2

100% Si 15.3% A1+ 84.7% Si 20.2% A1+ 79.8% Si 22.1% A1+ 77.9% Si 28.6% A1+ 71.4% Si 45.6% Cu + 54.4% Si 51.4% Cu + 48.6% Si

0.65-0.80 0.80 0.80 0.80 0.65 0.65 0.65

1450 1500 1500 1500 1500 1600 1600

5 60 60 60 5 30 30

systems has not been completed, in general we find that the minimum temperature at which spontaneous capillary infiltration occurs decreases as the silicon concentration of the melt increases. As one example, the samples SiC-Cul and SiC-Cu2 in Table 1 differ only in the composition of the ahoy. While the former, with 54.4% Si, exhibits complete infiltration under the conditions listed, the latter, which contains 48.6% Si, infiltrated only to a depth of 3-4 ram. (Except for SiC-Cu2, all the samples in Table 1 were completely infiltrated and, while the listed reaction times for silicon-metal samples are longer that for the pure silicon infiltration, we believe that they are actually fully reacted in a much shorter time.) After reactive infiltration and cooling, the excessive melt was removed from the surface of the samples. We accomplished this by using either a room temperature chemical etching process or a high temperature melt removal. For chemical etching, a mixture of concentrated acids, 70 wt.% of HF and 30 wt.% of HCI, was used. This mixture has been shown to give a fast etching rate for silicon [16,17], approaching 25/~m s-1. Changing the above ratio in either direction or diluting the acid concentration with water can considerably slow down the etching rate. In the high temperature removal process, the reacted composite is placed on a porous graphite plate or a bed of graphite powder and heated above the melting point of the alloy for a short time. The molten silicon reacts with the graphite leaving behind a smooth composite surface coated with a thin layer of silicon. 2.4. Simultaneous joining and reaction bonding The small dimensional changes which occur during reactive infiltration suggest that one can fabricate a complex-shape component from simpler parts machined separately. As a demonstration of simul-

taneous reaction bonding and joining, we machined a cylindrical pin, with an outer diameter closely matched to the inner diameter of a drilled hole in a separate preform. The assembled component was then reaction infiltrated with pure silicon (Fig. 5). In order to make the interface more easily observable, the two parts were prepared from different carbon densities (0.80 and 0.65 g cm-3). After infiltration, the component was sectioned and polished in order to observe the microstructure of the interface. 2.5. Evaluation of dimensional retention and surface finish Dimensional changes during preform infiltration and reaction were measured over sample gauge lengths of about 60 mm. The porous carbon preforms were machined in the shape of rectangular bars, and surface ground to ensure that opposing surfaces were fiat and parallel. The lengths of the samples were measured with a micrometer before and after reactive infiltration and removal of the excessive melt (by chemical etching). Also, a set of parallel reference lines was scribed on the preform surfaces prior to reaction and revealed after etching. The distances between lines were measured with a travelling stage microscope having a 2.5 a m displacement resolution. For surface finish retention studies, the surfaces of several bars were ground with silicon carbide abrasive papers: Buehler Carbimet grit 600, Struers # 1200, and Struers # 4000, corresponding to 14 pm, 14 pm and 5 pm average grain sizes respectively. The first two, although characterized by similar grain sizes, gave slightly different surface finishes probably because of different abrasive concentrations. The surface roughness before and after reactive infiltration was measured using a Dektak 3 surface profilometer (2.5 p m stylus, 30 mgf force). Samples were scanned over a 20 mm

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a

3. Evaluation of composites The microstructure and phase composition of composites were analyzed with optical microscopy, SEM, electron microprobe, and X-ray diffractometry (XRD). The ratio of normalized Si(100):SiC(100) X-ray peak heights was used as a measure of silicon content in SiC-Si composites. To quantify the relationship between peak heights and phase fraction, a calibration curve was prepared by using standards, made by mixing known amounts of fl-SiC (Johnson Matthey, 1 ktm average grain size) and silicon (Johnson Matthey, 100 mesh) powders. This method was also used for assessing the silicon content in metal-silicon-SiC composites. The metal or silicide content was then back calculated with the assumption that the metal-tosilicon ration in the composite is the same as that in the infiltrating alloy. The density of composites was measured using the Archimedes method, with methanol as the immersion fluid. The Young's modulus was measured by an acoustic resonance method using a GrindoSonic system according to ASTM standard C 1259 (courtesy of Dr. A n d r e Van Leuven from J.W. Lemmens, Inc., 10801 Pear Tree Lane, St. Louis, MO 63074). Flexural strength was measured at room temperature, on samples ground to 20 mm x 3 mm x 2 mm, using three-point bend test. Limited high temperature flexural measurements were also done in a vacuum furnace ( 10-4 Torr) at 1300 °C. Indentation fracture toughness K~c was measured using the Vickers indentation method [18]. A standard Vickers' indenter was mounted on an Instron testing machine and the sample was loaded with a constant cross-head speed of 0.05 mm min- 1, up to a maximum load ranging from 1.1 to 23 kgf. The maximum load was applied for about 15 s and removed at the same rate. The indentation sizes and crack lengths were measured from optical micrographs. For K~c calculations we used formulae given by Niihara et al. [19,20].

Fig. 5. Simultaneous joining and reactive infiltration. The component was made of two parts (stepped pin and bar) and was infiltrated with silicon after assembling: (a) after reactive infiltration; (b) after etching to remove excessive melt; (c) microstructure of the interface between the two parts. The thin, interfacial film of residual silicon is believed to form as a result of machining mismatch. The differing microstructures of the two parts are the result of differing preform densities. length. We used a low scanning speed with 2000 data points collected from each scan. Sample surface curvature or waviness due to machining, which is characterized by a much longer wavelength compared with the surface roughness, was arithmetically removed during data reduction.

4. Results and discussion 4.1. Microstructure and phase composition SiC-Si composites The microstructure of a typical SiC-Si composite is shown in Fig. 6, in which the darker phase is SiC and the brighter phase is the residual Si. In these composites, no porosity or residual carbon is visible through the reacted cross-sections of about 1 cm. The grain size of SiC is approximately 3 /zm, close to the carbon size scale in the starting preform. Although typically only the bottom 3 cm of a preform is

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Fig. 6. Microstructure of SiC-Si composite reactively infiltrated using glassy carbon preform of 0.80 g c m - 3 density. The darker phase is silicon carbide (86 vol.%) and the lighter phase is residual silicon ( 14 vol.%); viewed with optical microscopy.

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immersed in the melt, capillary infiltration allows the upper part of the preform to be also infiltrated throughout the entire preform length of 6 cm. No interface or difference in microstructure is visible between the lower and upper sections of such samples. In Fig. 7 we show typical X R D patterns for composites made from three different preform densities. Cubic SiC (fl-phase) and elemental Si are the only phases detected, and the relative intensities of SiC and Si lines systematically vary with the preform density. In Table 2 we list theoretical values for the volume fractions of SiC and Si, and experimental results determined by XRD. Depending on preform density, the volume fraction of silicon carbide is found to vary from 63% to 86% and the density of the composites from 2.89 to 3.06 g c m -3. The good agreement with calculated results shows that reaction is complete; indeed neither residual carbon nor porosity is seen in the microstructure.

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Microstructures of several SiC-AI composites of different preform densities and aluminum contents are shown in Fig. 8. The scale (grain size) of the microstructures appears to be very close to those seen for silicon-infiltrated samples (Fig. 6). After optimizing infiltration conditions, we did not observe any residual porosity or unreacted carbon in these composites. In Fig. 8 we can distinguish between three different phase regions: the darkest contrast phase is SiC, the intermediate gray phase is an A1-Si alloy, and the brightest phase is AI. Some amount of free silicon is always present, regardless of the aluminum content of the infiltrating alloy. However, the origin of this silicon is different from that in the SiC-Si composites, as discussed below. The XRD spectrum of sample SiC-AI2 is shown in Fig. 9. By comparison with Fig. 7(c) for the same preform density, we can see that the height of Si peaks relative to SiC has decreased, and peaks of metallic A1 have appeared. No detectable aluminum carbide peaks can be seen here, nor were any detected for any of the composites synthesized. The phase compositions of metal-infiltrated samples,

Table 2 Calculated and observed SiC and Si contents and composite densities in SiC-Si composites Preform density (g c m - 3)

0.63 _ 0.02 0.72 _+0.02 0.81 + 0.02

Theoretical

Measured

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Si (vol.%)

p (g c m - 3)

SiC (vol.%)

Si (vol.%)

p (g cm- 3)

65.5 74.9 84.2

34.5 25.1 15.8

2.92 3.02 3.06

62.9 + 2.0 78.0 + 2.0 86.1 + 2.0

37.1 + 2.0 22.0 + 2.0 13.9 + 2.0

2.89 + 0.02 2.98 + 0.02 3.05 + 0.02

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Table 3 Phase compositions of silicon-metal infiltrated samples calculated on the basis of silicon contents measurements

Fig. 8. SiC-A1 composites reactively infiltrated at 1500 °C using Si-AI alloys, viewed with optical microscopy: (a) carbon preform of 0.80 g c m -3 density infiltrated with 15 mol.% A1 + 85 mol.% Si alloy; (b) carbon preform of 0.80 g cm -3 density infiltrated with 20 mol.% AI + 80 mol.% Si alloy; (c) carbon preform of 0.65 g cm -3 density infiltrated with 29 mol.% AI + 71 mol.% Si alloy.

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Phase composition (vol.%)

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determined using XRD peak heights, are fisted in Table 3. The SiC volume fraction is slightly smaller than expected from the preform density, which we attribute to measurement error. The aluminum volume fraction varies from 13% to 24%. We believe that the composite formation process is as follows. During reactive infiltration, the Si-AI liquid penetrates the preform and Si selectively reacts with carbon to form SiC, leaving as the secondary phase an aluminum-enriched melt. The compositions have been chosen so that, at infiltration temperature, the overall composition falls into a two-phase region of the AI-Si-C ternary phase diagram, where SiC and an AI-Si liquid are in equilibrium. This A1-Si liquid is of near-eutectic composition. On cooling, the liquid crystallizes, and may first solidify some primary silicon or aluminum, depending on whether it is of hyper- or hypoeutectic composition. The appearance of primary A1 crystals in the microstructures of Al-rich samples

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SiC-A13 and SiC-AI4 (Figs. 8(b) and 8(c)) is consistent with hypoeutectic residual liquid. The majority of the liquid then undergoes eutectic solidification. Analytical electron microscopy revealed finely divided silicon-aluminum microstructures characteristic of eutectic solidification. Microstructures of samples infiltrated with Si-Cu alloys, as shown in Fig. 10, also consist of three phases: SiC, Si, and a bright contrast copper-rich phase. The copper-rich phase was analyzed using electron microprobe microanalysis and found to contain 70 + 7 at.% of Cu and 30 + 7 at.% of Si. (The relatively high error in the compositional analysis results from the small size of the analyzed particles.) This composition is close to that of the r/' copper silicide phase. XRD spectra (Fig. 11) do not show the existence of metallic copper in these composites and show unidentified peaks which we attribute to copper silicide, even though an exact match to patterns in the JCPDS X-ray database was not found. Unlike the Si-C-AI system, the Si-C-Cu phase diagram is not known with confidence, especially at the infiltration temperature used here. Thus the equilibrium phases at the reaction temperature cannot be independently confirmed. However, based on the phases observed in the reacted composites we can provide the following qualitative description of the phase formation process. It is clear that the SiC phase is first formed, causing the remaining melt to be Cu rich. Since our reaction temperature is above the melting points of Si and the copper silicides, it is likely that the SiC coexists in equilibrium with a single-phase Cu-Si liquid. As in the AI-Si case, this liquid solidifies

as essentially a binary system. The presence of silicon in the final composite, as well as the composition of the Cu-rich phase, shows that the secondary melt composition is on the silicon-rich side of the eutectic in the Cu-Si binary system [15]. Additional details of the compositional design of these composites and the phase formation process are to be published elsewhere. Based on observations to date, it appears that in both the Si-C-AI and the Si-C-Cu systems, the SiC phase forms rapidly and coexists at reaction temperatures ( 1500-1600 °C) with a single-phase A1-Si or Cu-Si liquid, which fills the pore space of the composite. Thus the secondary phase content of the composite is determined by the solidification of this liquid. It seems possible to control further the secondary phases by varying overall composition and temperature in this process; of particular interest is the degree to which the phase(s) can be enriched in A1 or Cu. Additional work to establish the compositional limits for the residual phases in these two systems is in progress. 4.2. Dimensional retention In Table 4 we show typical results for the linear shrinkage A L / L between a machined carbon preform and the reacted composite, measured with a resolution better than 0.01%. While there is a noticeable trend of increasing shrinkage with decreasing carbon density, in no case does the shrinkage exceed 0.5%. For the highest density of preforms, several experiments showed that the dimensional change remains less than 0.1%. This was observed for composites of all of the compositions studied in this work.

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Rigaku300 50kV;200 mA; 10°/min.,sampling0.05 Al*-signalfromAl sample holder Fig. 10. Reactively infiltrated SiC-Cu composite, prepared using carbon preform of 0.65 g c m -3 density, and infiltrated with 45 mol.% Cu + 55 mol.% Si alloy at 1600 °C, viewed with optical microscopy.

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Table 4 Typical linear dimensionalchanges of SiC-Si compositesafter reactive infiltration Preform density (g cm- 3)

Initial length (mm)

Final length (mm)

Change of linear dimension (%)

0.65 0.73 0.80

60.45 61.05 62.57

60.21 60.89 62.53

- 0.41 - 0.26 - 0.06

In order to understand how this high level of dimensional retention is achieved, it is useful to consider first the molar volume change on reacting glassy carbon to silicon carbide. For glassy carbon of 1.5 g cm -3 density (8 cm 3 mol-1) there is a 50% increase in molar volume on forming SiC (12.5 cm 3 mol-l). Secondly, we must take into account the fact that the dimensional change measured at room temperature also includes the thermal contraction of the composite on cooling from the reaction temperature. The thermal expansion coefficient of siliconized silicon carbide containing 25 vol.% of silicon is estimated using the linear rule of mixtures to be about 5.0 × 10-6 °C- l (in the temperature range 25-1500 °C). Accordingly, our reacted composite should decrease in length by about 0.75% during cooling from the infiltration temperature. (Of course, the expansion of carbon on heating should, strictly speaking, also be taken into account, but it is a small correction.) Correcting for thermal contraction, we thus find that there has actually been a slight expansion of the preform during reactive infiltration at high temperature. However, this expansion is much less than the 50% expansion expected for direct transformation of a carbon network to SiC. This is consistent with the solution precipitation mechanism of reaction between liquid Si and C [6-8], in which we expect that the carbon network exists as a solid skeleton until sufficient SiC has formed to become the interconnected solid network. The high degree of dimensional retention suggests that reactive infiltration can truly be a net-shape process. These dimensional changes are far less than those in any sintering-based process, where typical shrinkages are about 30% for fully densified materials. As a demonstration of net-shape capability, we show in Fig. 4 an example of a reaction-infiltrated bolt (1/4 in × 20 threads in- 1), which remains closely matched to the steel hexagonal nut after reaction and removal of the excess melt. Another example of a net-shape processed prototype component is shown in Fig. 4, in the form of an internal combustion engine valve.

4.3. Infiltration of assembled component In Fig. 5 we show the two-part component consisting of cylindrical pin mated to a square bar, after infiltration and after removal of the excess melt. As expected, the shape of the part was accurately retained. The microstructure of the interface is shown in Fig. 5(c). The difference in microstructure between the two parts is due to the difference in preform density. A thin interfacial film is visible, the thickness of which seems to depend on the fit between the assembled parts (accuracy of machining), since it was not visible at other regions along the interface. This demonstration shows that, in principle, complex parts can be simultaneously joined and reacted, and also that parts with dissimilar microstructures and properties, or possibly graded microstructures, can be incorporated as well.

4.4. Surface finish retention In Figs. 12 and 13 we show the measurement of surface roughness before and after infiltration and etching for two preform densities, 0.65 g cm -3 and 0.80 g cm-3 respectively. Note the differences between the vertical scale (micrometers) and horizontal scale (millimeters). As increasingly finer grit sizes are used for surface preparation, the surface roughness R a of the carbon preforms becomes dictated by the pore structure, and is of the order of 1.0-1.3 /~m. This surface roughness is retained after processing and is of the order of 1.2-1.5 /~m. Microscopic observation shows that the roughness after processing is limited by the fact that residual silicon is etched from pores near the surface of the sample and reflects the scale of the SiC microstructure. This level of surface roughness may be suitable for some applications, while others would require further finishing.

4.5. Physical properties Some physical properties of the processed composites are listed in Table 5. The measured Young's modulus E of the SiC-Si composites was 350 GPa. Using a value of E = 4 6 6 GPa for SiC based on measurements of chemical vapor deposition grown, very high purity and 100% dense SiC [21], and E = 113 GPa for Si, the linear rule of mixtures gives E = 395 GPa for a composite with 20 vol.% of silicon. However, a negative deviation from the linear rule of mixtures is typical of composite mixtures [22]. As expected, the presence of aluminum ( E = 7 0 - 8 0 GPa) in the SiC-A1-Si composites decreases slightly the Young's modulus relative to that of SiC-Si. Room temperature flexural strengths were in the range 400-600 MPa for SiC-Si composites, slightly

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before infiltration

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Ra=3.51 prn

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141

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Ra=1.52 pm

Struers # 4000 (5 lam) 0

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0

5

10

Fig. 12. Surface r o u g h n e s s profile o f t h e 0.65 g c m -3 glassy c a r b o n p r e f o r m m e a s u r e d b e f o r e and after reactive infiltration and chemical etching. Results for differing initial surface finish are shown.

lower (370-390 MPa) for SiC-A1-Si composites, and about the same (510 MPa) for SiC-Cu-Si composites. Each of these values is the average of at least three tested specimens. Values of about 550 MPa have been observed in limited high temperature tests of the SiC-Si composites (1300 °C in vacuum). In general, these strengths are a significant improvement over those of conventional RBSCs, which can be attributed to the finer-grained microstructures and possibly a more highly interconnected carbide phase in the present composites. The Vickers microhardness of the SiC-Si composites was in the range 17-18 GPa, varying slightly with the Si content, and decreases, as expected, in the SiC-metal system on addition of the softer metal phase. The indentation fracture toughness of the

SiC-Si composites is similar to that of conventional RBSC (4.0 MPa mU2). The fracture toughness of the SiC-metal composites is modestly better (5.0-5.7 MPa m]/2). These values seem low considering the significant metal fraction (Table 3), which we can attribute to the fact that the majority of the residual metal is present in a phase or microstructure which is not of high ductility. In the SiC-A1-Si composites, most of the aluminum is the result of eutectic solidification, which is known to yield low toughness material (by metallurgical standards). In the SiC-Cu-Si composites, the copper is present primarily as a copper silicide, also not expected to be of high toughness. It is clear that, if higher fracture toughnesses are to be achieved, further control of the secondary phase compositions is necessary. In large part this involves a better understanding

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before infiltration

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Ra=2.05 pm

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Fig. 13. Surface roughness profile of the 0.80 g cm -3 glassy carbon preform measured before and after reactive infiltration and chemical etching. The results for differing initial surface finish are shown.

Table 5 Mechanical properties of SiC-Si and SiC-Si-metal composites Sample

Density (g c m - 3)

Young's modulus (GPa)

Vickers' hardness (GPa)

Indentation Kic (MPa m ]/2)

Flexural strength (MPa)

SiC-Si SiC-All SiC-A12 SiC-A13 SiC-AI4 SiC-Cu 1 SiC-Cu2

2.85- 3.11 3.05 2.99 3.03 2.93 3.71 Not measured

350 323 300 341 299 Not measured Not measured

17-18 17.84 ( + 1.05) 17.71 ( + 0.97) 14.78 ( + 1.91 ) 10.85 ( _+ 1.16) 12.60 ( _ 1.80) 13.90 ( + 1.20)

3.5-4.5 4.5 ( _+0.4) 4.0 ( _+0.3) 5.7 ( _+ 1.9) 5.5 ( -+ 0.8) 5.0 ( _+0.4) 5.3 ( + 0.5)

400-600 390 Not measured Not measured 370 510 Not measured

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of the high temperature phase equilibria and manipulation of thermal history during processing and is the subject of ongoing work. 5. Summary Silicon carbide composites in which both the major and the minor phase are interpenetrating, and in which the minor phase is silicon or metal-silicon alloys, have been fabricated by pressureless reactive infiltration of porous glassy carbon preforms. The composites are characterized by homogeneous, fine-grain microstructure, in which the SiC content is controlled by varying the carbon density of the preform. By alloying A1 or Cu melts with Si, we are able to achieve pressureless infiltration of these otherwise non-wetting metals and to introduce metal-rich secondary phases. The Young's modulus and Vickers hardness of the composites are comparable with those of previous siliconized silicon carbides; the flexural strengths and fracture toughnesses are improved. It is demonstrated that this reactive infiltration process is capable of producing net-shape components with a linear dimension change of less than 0.1%. The surface finish of the components after reaction and removal of excess melt varies with the initial preform finish, but can be better than Ra = 1.5/,m. This is comparable with that achieved after conventional lapping. The simultaneous joining and reactive infiltration of complex parts made by assembling simpler multipart preforms has also been demosntrated.

Acknowledgments We thank Dr. Kristen R Constant for assistance with preform preparation and Ling Liao for assistance with indentation measurements. This work was supported by the Office of Naval Research, Grants N00014-90-J-1999 and N0001494-1-0790, Dr. Steven G. Fishman, Program Manager. References [1] P. Popper, The preparation of dense self bonded silicon carbide, in Special Ceramics, Heywood, London, 1960, p. 209. [2] C.W. Forrest, P. Kennedy and J.V. Shennan, The fabrication and properties of self-bonded silicon carbide bodies, in P. Popper (ed.), Special Ceramics, Vol. 5, British Ceramic Research Association, Stoke-on-Trent, 1972, p. 99. [3] W.B. Hillig, R.L. Mehan, C.R. Morelock, V.J. DeCarlo and W. Laskow, Silicon/silicon carbide composites, Am. Ceram. Soc. Bull., 54(12)(1975) 1054-1056.

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[4] E.E. Hucke, Composite bodies comprising a continuous framework and an impregnated metallic material and methods of their production, US Patent 3,235,346, February 15, 1966. [5] E.E. Hucke, Process development for silicon carbide based structural ceramics, A M M R C Tech. Rep. TR 83-5, January 1983 (US Army Materials and Mechanics Research Center). [6] Y.-M. Chiang, R.R Messner, C.D. Terwilliger and D.R. Behrendt, Reaction-formed silicon carbide, Mater. Sci. Eng. A, 144(1991) 63-74. [7] J.N. Ness and T.F. Page, Microstructural evolution in reaction-bonded silicon carbide, J. Mater. Sci., 21 (4) (1986) 1377-1397. [8] R. Pampuch, E. Walasek and J. Bialoskbrski, Reaction mechanism in carbon-liquid silicon systems at elevated temperatures, Ceram. Int., 12 (2) ( 1986 ) 99-106. [9] W.-Y. Lin, J.-Y. Hsu, Y. Berta and R.E Speyer, Combustion gas corrosion resistance of heat exchange materials and refractories for glass furnaces at high temperatures: part I. Silicon carbide and molybdenum silicide, Am. Ceram. Soc. Bull., 73(2)(1994)72-78. [10] R.P. Messner and Y.-M. Chiang, Processing of reactionbonded silicon carbide without residual silicon phase, Ceram. Eng. Sci. Proc., 9(7-8)(1988) 1052-1060. [11] R.R Messner and Y.-M. Chiang, Liquid-phase reactionbonding of silicon carbide using alloyed silicon-molybdenum melts, J. Am. Ceram. Soc., 73 (5) (1990) 1193-1200. [12] F. Delannay, L. Froyen and A. Deruyttere, Review. The wetting of solids by molten metals and its relation to the preparation of metal-matrix composites, J. Mater. Sci., 22 (1)(1987) 1-16. [13] C. Zweben, Metal-matrix composites for electronic packaging, J. Met., 44(7)(1992) 15-23. [14] Y.-L. Shen, A. Needleman, and S. Suresh,d Coefficients of thermal expansion of metal-matrix composites for electronic packaging, Metall. Met. Trans. A, 25 (4) (1994) 839-849. [15] T.B. Massalski (ed.), Binary Alloy Phase Diagrams, ASM International, Materials Park, OH, 2nd edn., 1990. [16] D.R. Turner, On the mechanism of chemically etching germanium and silicon, J. Electrochem. Soc., 107 (10) (1960) 810-816. [17] H. Robbins and B. Schwartz, Chemical etching of silicon I. The system HF, HNO3, and H2, J. Electrochem. Soc., 106 (6) (1959) 505-508. [18] B.R. Lawn and D.B. Marshall, Indentation fracture and strength degradation in ceramics, in R.C. Bradt, D.P.H. Hasselman and EF. Lange (eds.), Fracture Mechanics of Ceramics, Vol. 3, Flaws and testing, Plenum, New York, 1978, p. 205. [19] N. Niihara, A fracture mechanics analysis of indentationinduced Palmqvist crack in ceramics, J. Mater. Sci. Lett., 2 (5)(1983)221-223. [20] K. Niihara, R. Morena and D.P.H. Hasselman, Evaluation of Kit of brittle solids by the indentation method with low crack-to-indent ratios, J. Muter. Sci. Lett., 1 (1) (1982) 13-16. [21] Manufacturing innovations, CVD scaled up for commercial production of bulk SiC, Am. Ceram. Soc. Bull., 72 (3) (1993) 74-78. [22] Z. Hashin and S. Shtrikman, A variational approach to the theory of the elastic behavior of multiphase materials, J. Mech. Phys. Solids, 11 (1963) 127-140.